Edge formability in metallic alloys

ABSTRACT

This disclosure is directed at methods for mechanical property improvement in a metallic alloy that has undergone one or more mechanical property losses as a consequence of forming an edge, such as in the formation of an internal hole or an external edge. Methods are disclosed that provide the ability to improve mechanical properties of metallic alloys that have been formed with one or more edges placed in the metallic alloy by a variety of methods which may otherwise serve as a limiting factor for industrial applications.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application is a continuation-in-part of U.S. patent applicationSer. No. 15/094,554 filed Apr. 8, 2016, which claims the benefit of U.S.Provisional Patent Application Ser. No. 62/146,048 filed on Apr. 10,2015 and U.S. Provisional Patent Application Ser. No. 62/257,070 filedon Nov. 18, 2015, which is fully incorporated herein by reference.

FIELD OF INVENTION

This disclosure relates to methods for mechanical property improvementin a metallic alloy that has undergone one or more mechanical propertylosses as a consequence of shearing, such as in the formation of asheared edge portion or a punched hole. More specifically, methods aredisclosed that provide the ability to improve mechanical properties ofmetallic alloys that have been formed with one or more sheared edgeswhich may otherwise serve as a limiting factor for industrialapplications.

BACKGROUND

From ancient tools to modern skyscrapers and automobiles, steel hasdriven human innovation for hundreds of years. Abundant in the Earth'scrust, iron and its associated alloys have provided humanity withsolutions to many daunting developmental barriers. From humblebeginnings, steel development has progressed considerably within thepast two centuries, with new varieties of steel becoming available everyfew years. These steel alloys can be broken up into three classes basedupon measured properties, in particular maximum tensile strain andtensile stress prior to failure. These three classes are: Low StrengthSteels (LSS), High Strength Steels (HSS), and Advanced High StrengthSteels (AHSS). Low Strength Steels (LSS) are generally classified asexhibiting ultimate tensile strengths less than 270 MPa and include suchtypes as interstitial free and mild steels. High-Strength Steels (HSS)are classified as exhibiting ultimate tensile strengths from 270 to 700MPa and include such types as high strength low alloy, high strengthinterstitial free and bake hardenable steels. Advanced High-StrengthSteels (AHSS) steels are classified by ultimate tensile strengthsgreater than 700 MPa and include such types as Martensitic steels (MS),Dual Phase (DP) steels, Transformation Induced Plasticity (TRIP) steels,and Complex Phase (CP) steels. As the strength level increases the trendin maximum tensile elongation (ductility) of the steel is negative, withdecreasing elongation at high ultimate tensile strengths. For example,tensile elongation of LSS, HSS and AHSS ranges from 25% to 55%, 10% to45%, and 4% to 30%, respectively.

Production of steel continues to increase, with a current US productionaround 100 million tons per year with an estimated value of $75 billion.Steel utilization in vehicles is also high, with advanced high strengthsteels (AHSS) currently at 17% and forecast to grow by 300% in thecoming years [American Iron and Steel Institute. (2013). Profile 2013.Washington, D.C.]. With current market trends and governmentalregulations pushing towards higher efficiency in vehicles, AHSS areincreasingly being pursued for their ability to provide high strength tomass ratio. The high strength of AHSS allows for a designer to reducethe thickness of a finished part while still maintaining comparable orimproved mechanical properties. In reducing the thickness of a part,less mass is needed to attain the same or better mechanical propertiesfor the vehicle thereby improving vehicle fuel efficiency. This allowsthe designer to improve the fuel efficiency of a vehicle while notcompromising on safety.

One key attribute for next generation steels is formability. Formabilityis the ability of a material to be made into a particular geometrywithout cracking, rupturing or otherwise undergoing failure. Highformability steel provides benefit to a part designer by allowing forthe creation of more complex part geometries allowing for reduction inweight. Formability may be further broken into two distinct forms: edgeformability and bulk formability. Edge formability is the ability for anedge to be formed into a certain shape. Edges on materials are createdthrough a variety of methods in industrial processes, including but notlimited to punching, shearing, piercing, stamping, perforating, cutting,or cropping. Furthermore, the devices used to create these edges are asdiverse as the methods, including but not limited to various types ofmechanical presses, hydraulic presses, and/or electromagnetic presses.Depending upon the application and material undergoing the operation,the range of speeds for edge creation is also widely varying, withspeeds as low as 0.25 mm/s and as high as 3700 mm/s. The wide variety ofedge forming methods, devices, and speeds results in a myriad ofdifferent edge conditions in use commercially today.

Edges, being free surfaces, are dominated by defects such as cracks orstructural changes in the sheet resulting from the creation of the sheetedge. These defects adversely affect the edge formability during formingoperations, leading to a decrease in effective ductility at the edge.Bulk formability on the other hand is dominated by the intrinsicductility, structure, and associated stress state of the metal duringthe forming operation. Bulk formability is affected primarily byavailable deformation mechanisms such as dislocations, twinning, andphase transformations. Bulk formability is maximized when theseavailable deformation mechanisms are saturated within the material, withimproved bulk formability resulting from an increased number andavailability of these mechanisms.

Edge formability can be measured through hole expansion measurements,whereby a hole is made in a sheet and that hole is expanded by means ofa conical punch. Previous studies have shown that conventional AHSSmaterials suffer from reduced edge formability compared with other LSSand HSS when measured by hole expansion [M. S. Billur, T. Altan,“Challenges in forming advanced high strength steels”, Proceedings ofNew Developments in Sheet Metal Forming, pp. 285-304, 2012]. Forexample, Dual Phase (DP) steels with ultimate tensile strength of 780MPa achieve less than 20% hole expansion, whereas Interstitial Freesteels (IF) with ultimate tensile strength of approximately 400 MPaachieve around 100% hole expansion ratio. This reduced edge formabilitycomplicates adoption of AHSS in automotive applications, despitepossessing desirable bulk formability.

SUMMARY

A method for improving one or more mechanical properties in a metallicalloy that has undergone a mechanical property loss as a consequence ofthe formation of one or more sheared edges comprising:

-   -   a. supplying a metal alloy comprising at least 50 atomic % iron        and at least four or more elements selected from Si, Mn, B, Cr,        Ni, Cu or C and melting said alloy and cooling at a rate of ≤250        K/s or solidifying to a thickness of ≥2.0 mm up to 500 mm and        forming an alloy having a T_(m) and matrix grains of 2 μm to        10,000 μm;    -   b. heating said alloy to a temperature of ≥700° C. and below the        T_(m) of said alloy and at a strain rate of 10⁻⁶ to 10⁴ and        reducing said thickness of said alloy and providing a first        resulting alloy having an ultimate tensile strength of 921 MPa        to 1413 MPa;    -   c. stressing said first resulting alloy and providing a second        resulting alloy having an ultimate tensile strength of 1356 MPa        to 1831 MPa and an elongation of 1.6% to 32.8%;    -   d. heating said second resulting alloy to a temperature below        T_(m) and forming a third resulting alloy having matrix grains        of 0.5 μm to 50 μm and having an elongation (E₁);    -   e. shearing said alloy and forming one or more sheared edges        wherein said alloy's elongation is reduced to a value of E₂        wherein E₂=(0.57 to 0.05) (E₁)    -   f. reheating said alloy with said one or more sheared edges        wherein said alloy's reduced elongation observed in step (d) is        restored to a level having an elongation E₃=(0.48 to 1.21)(E₁).

The present disclosure also relates to a method for improving the holeexpansion ratio in a metallic alloy that had undergone a hole expansionratio loss as a consequence of forming a hole with a sheared edgecomprising:

-   -   a. supplying a metal alloy comprising at least 50 atomic % iron        and at least four or more elements selected from Si, Mn, B, Cr,        Ni, Cu or C and melting said alloy and cooling at a rate of ≤250        K/s or solidifying to a thickness of ≥2.0 mm up to 500 mm and        forming an alloy having a T_(m) and matrix grains of 2 μm to        10,000 μm;    -   b. heating said alloy to a temperature of 700° C. and below the        T_(m) of said alloy and at a strain rate of 10⁻⁶ to 10⁴ and        reducing said thickness of said alloy and providing a first        resulting alloy having an ultimate tensile strength of 921 MPa        to 1413 MPa and an elongation of 12.0% to 77.7%;    -   c. stressing said first resulting alloy and providing a second        resulting alloy having an ultimate tensile strength of 1356 MPa        to 1831 MPa and an elongation of 1.6% to 32.8%;    -   d. heating said second resulting alloy to a temperature of at        least 650° C. and below Tm and forming a third resulting alloy        having matrix grains of 0.5 μm to 50 μm and forming a hole        therein with shearing wherein said hole has a sheared edge and        has a first hole expansion ratio (HER₁);    -   e. heating said alloy with said hole and associated HER₁ wherein        said alloy indicates a second hole expansion ratio (HER₂)        wherein HER₂≥HER₁.

The present invention also relates to method for improving the holeexpansion ratio in a metallic alloy that had undergone a hole expansionratio loss as a consequence of forming a hole with a sheared edgecomprising:

-   -   a. supplying a metal alloy comprising at least 50 atomic % iron        and at least four or more elements selected from Si, Mn, B, Cr,        Ni, Cu or C and melting said alloy and cooling at a rate of ≤250        K/s or solidifying to a thickness of ≥2.0 mm up to 500 mm and        forming an alloy having a Tm and matrix grains of 2 μm to 10,000        μm;    -   b. heating said alloy to a temperature of ≥700° C. and below the        Tm of said alloy and at a strain rate of 10⁻⁶ to 10⁴ and        reducing said thickness of said alloy and providing a first        resulting alloy having an ultimate tensile strength of 921 MPa        to 1413 MPa and an elongation of 12.0% to 77.7%;    -   c. stressing said first resulting alloy and providing a second        resulting alloy having an ultimate tensile strength of 1356 MPa        to 1831 MPa and an elongation of 1.6% to 32.8%;    -   d. heating said second resulting alloy to a temperature of at        least 650° C. and below Tm and forming a third resulting alloy        having matrix grains of 0.5 μm to 50 μm wherein said alloy is        characterized as having a first hole expansion ratio (HER₁) of        30 to-130% for a hole formed therein without shearing;    -   e. forming a hole in said second resulting alloy wherein said        hole is formed with shearing and indicates a second hole        expansion ratio (HER₂) wherein HER₂=(0.01 to 0.30)(HER₁);    -   f. heating said alloy wherein HER₂ recovers to a value        HER₃=(0.60 to 1.0) HER₁.

The present invention also relates to a method for punching one or moreholes in a metallic alloy comprising:

-   -   a. supplying a metal alloy comprising at least 50 atomic % iron        and at least four or more elements selected from Si, Mn, B, Cr,        Ni, Cu or C and melting said alloy and cooling at a rate of ≤250        K/s or solidifying to a thickness of ≥2.0 mm up to 500 mm and        forming an alloy having a Tm and matrix grains of 2 μm to 10,000        μm;    -   b. heating said alloy to a temperature of 700° C. and below the        Tm of said alloy and at a strain rate of 10⁻⁶ to 10⁴ and        reducing said thickness of said alloy and providing a first        resulting alloy having an ultimate tensile strength of 921 MPa        to 1413 MPa and an elongation of 12.0% to 77.7%;    -   c. stressing said first resulting alloy and providing a second        resulting alloy having an ultimate tensile strength of 1356 MPa        to 1831 MPa and an elongation of 1.6% to 32.8%;    -   d. heating said second resulting alloy to a temperature of at        least 650° C. and below Tm and forming a third resulting alloy        having matrix grains of 0.5 μm to 50 μm and having an elongation        (E₁);    -   e. punching a hole in said alloy at a punch speed of greater        than or equal to 10 mm/second wherein said punched hole        indicates a hole expansion ratio of greater than or equal to        10%.

The present invention also relates to a method for expanding an edge inan alloy

-   -   a. supplying a metal alloy comprising at least 50 atomic % iron        and at least four or more elements selected from Si, Mn, B, Cr,        Ni, Cu or C and melting said alloy and cooling at a rate of ≤250        K/s or solidifying to a thickness of ≥2.0 mm up to 500 mm and        forming an alloy having a Tm;    -   b. heating said alloy to a temperature of 700° C. and below the        Tm of said alloy and at a strain rate of 10⁻⁶ to 10⁴ and        reducing said thickness of said alloy and providing a first        resulting alloy having an ultimate tensile strength of 921 MPa        to 1413 MPa and an elongation of 12.0% to 77.7%;    -   c. stressing said first resulting alloy and providing a second        resulting alloy having an ultimate tensile strength of 1356 MPa        to 1831 MPa and an elongation of 1.6% to 32.8%;    -   d. heating said second resulting alloy to a temperature of below        Tm and forming a third resulting alloy having an elongation of        6.6% to 86.7%;    -   e. forming an edge in said resulting alloy and expanding said        edge at a speed of greater than or equal to 5 mm/min.

The present invention also relates to a method for expanding the edge ofan alloy comprising:

-   -   supplying a metal alloy comprising at least 50 atomic % iron and        at least four or more elements selected from Si, Mn, B, Cr, Ni,        Cu or C, wherein said alloy has an ultimate tensile strength of        799 MPa to 1683 MPa and an elongation of 6.6 to 86.7%;    -   forming an edge in said alloy;    -   expanding said edge in said alloy at a speed of greater than or        equal to 5 mm/min.

BRIEF DESCRIPTION OF THE DRAWINGS

The detailed description below may be better understood with referenceto the accompanying FIG.s which are provided for illustrative purposesand are not to be considered as limiting any aspect of this invention.

FIG. 1A Structural pathway for the formation of High Strength NanomodalStructure and associated mechanisms.

FIG. 1B Structural pathway for the formation of Recrystallized ModalStructure and Refined High Strength Nanomodal Structure and associatedmechanisms.

FIG. 2 Structural pathway toward developing Refined High StrengthNanomodal Structure which is tied to industrial processing steps.

FIG. 3 Images of laboratory cast 50 mm slabs from: a) Alloy 9 and b)Alloy 12.

FIG. 4 Images of hot rolled sheet after laboratory casting from: a)Alloy 9 and b) Alloy 12.

FIG. 5 Images of cold rolled sheet after laboratory casting and hotrolling from: a) Alloy 9 and b) Alloy 12.

FIG. 6 Microstructure of solidified Alloy 1 cast at 50 mm thickness: a)Backscattered SEM micrograph showing the dendritic nature of the ModalStructure in the as-cast state, b) Bright-field TEM micrograph showingthe details in the matrix grains, c) Bright-field TEM with selectedelectron diffraction exhibiting the ferrite phase in the ModalStructure.

FIG. 7 X-ray diffraction pattern for the Modal Structure in Alloy 1alloy after solidification: a) Experimental data, b) Rietveld refinementanalysis.

FIG. 8 Microstructure of Alloy 1 after hot rolling to 1.7 mm thickness:a) Backscattered SEM micrograph showing the homogenized and refinedNanomodal Structure, b) Bright-field TEM micrograph showing the detailsin the matrix grains.

FIG. 9 X-ray diffraction pattern for the Nanomodal Structure in Alloy 1after hot rolling: a) Experimental data, b) Rietveld refinementanalysis.

FIG. 10 Microstructure of Alloy 1 after cold rolling to 1.2 mmthickness: a) Backscattered SEM micrograph showing the High StrengthNanomodal Structure after cold rolling, b) Bright-field TEM micrographshowing the details in the matrix grains.

FIG. 11 X-ray diffraction pattern for the High Strength NanomodalStructure in Alloy 1 after cold rolling: a) Experimental data, b)Rietveld refinement analysis.

FIG. 12 Bright-field TEM micrographs of microstructure in Alloy 1 afterhot rolling, cold rolling and annealing at 850° C. for 5 min exhibitingthe Recrystallized Modal Structure: a) Low magnification image, b) Highmagnification image with selected electron diffraction pattern showingcrystal structure of austenite phase.

FIG. 13 Backscattered SEM micrographs of microstructure in Alloy 1 afterhot rolling, cold rolling and annealing at 850° C. for 5 min exhibitingthe Recrystallized Modal Structure: a) Low magnification image, b) Highmagnification image.

FIG. 14 X-ray diffraction pattern for the Recrystallized Modal Structurein Alloy 1 after annealing: a) Experimental data, b) Rietveld refinementanalysis.

FIG. 15 Bright-field TEM micrographs of microstructure in Alloy 1showing Refined High Strength Nanomodal Structure (MixedMicroconstituent Structure) formed after tensile deformation: a) Largegrains of untransformed structure and transformed “pockets” with refinedgrains; b) Refined structure within a “pocket”.

FIG. 16 Backscattered SEM micrographs of microstructure in Alloy 1showing Refined High Strength Nanomodal Structure (MixedMicroconstituent Structure): a) Low magnification image, b) Highmagnification image.

FIG. 17 X-ray diffraction pattern for Refined High Strength NanomodalStructure in Alloy 1 after cold deformation: a) Experimental data, b)Rietveld refinement analysis.

FIG. 18 Microstructure of solidified Alloy 2 cast at 50 mm thickness: a)Backscattered SEM micrograph showing the dendritic nature of the ModalStructure in the as-cast state, b) Bright-field TEM micrograph showingthe details in the matrix grains.

FIG. 19 X-ray diffraction pattern for the Modal Structure in Alloy 2after solidification: a) Experimental data, b) Rietveld refinementanalysis.

FIG. 20 Microstructure of Alloy 2 after hot rolling to 1.7 mm thickness:a) Backscattered SEM micrograph showing the homogenized and refinedNanomodal Structure, b) Bright-field TEM micrograph showing the detailsin the matrix grains.

FIG. 21 X-ray diffraction pattern for the Nanomodal Structure in Alloy 2after hot rolling: a) Experimental data, b) Rietveld refinementanalysis.

FIG. 22 Microstructure of Alloy 2 after cold rolling to 1.2 mmthickness: a) Backscattered SEM micrograph showing the High StrengthNanomodal Structure after cold rolling, b) Bright-field TEM micrographshowing the details in the matrix grains.

FIG. 23 X-ray diffraction pattern for the High Strength NanomodalStructure in Alloy 2 after cold rolling: a) Experimental data, b)Rietveld refinement analysis.

FIG. 24 Bright-field TEM micrographs of microstructure in Alloy 2 afterhot rolling, cold rolling and annealing at 850° C. for 10 min exhibitingthe Recrystallized Modal Structure: a) Low magnification image, b) Highmagnification image with selected electron diffraction pattern showingcrystal structure of austenite phase.

FIG. 25 Backscattered SEM micrographs of microstructure in Alloy 2 afterhot rolling, cold rolling and annealing at 850° C. for 10 min exhibitingthe Recrystallized Modal Structure: a) Low magnification image, b) Highmagnification image.

FIG. 26 X-ray diffraction pattern for the Recrystallized Modal Structurein Alloy 2 after annealing: a) Experimental data, b) Rietveld refinementanalysis.

FIG. 27 Microstructure in Alloy 2 showing Refined High StrengthNanomodal Structure (Mixed Microconstituent Structure) formed aftertensile deformation: a) Bright-field TEM micrographs of transformed“pockets” with refined grains; b) Back-scattered SEM micrograph of themicrostructure.

FIG. 28 X-ray diffraction pattern for Refined High Strength NanomodalStructure in Alloy 2 after cold deformation: a) Experimental data, b)Rietveld refinement analysis.

FIG. 29 Tensile properties of Alloy 1 at various stages of laboratoryprocessing.

FIG. 30 Tensile results for Alloy 13 at various stages of laboratoryprocessing.

FIG. 31 Tensile results for Alloy 17 at various stages of laboratoryprocessing.

FIG. 32 Tensile properties of the sheet in hot rolled state and aftereach step of cold rolling/annealing cycles demonstrating full propertyreversibility at each cycle in: a) Alloy, b) Alloy 2.

FIG. 33 A bend test schematic showing a bending device with two supportsand a former (International Organization for Standardization, 2005).

FIG. 34 Images of bend testing samples from Alloy 1 tested to 180°: a)Picture of a full set of samples tested to 180° without cracking, and b)A close-up view of the bend of a tested sample.

FIG. 35 a) Tensile test results of the punched and EDM cut specimensfrom selected alloys demonstrating property decrease due to punched edgedamage, b) Tensile curves of the selected alloys for EDM cut specimens.

FIG. 36 SEM images of the specimen edges in Alloy 1 after a) EDM cuttingand b) Punching.

FIG. 37 SEM images of the microstructure near the edge in Alloy 1: a)EDM cut specimens and b) Punched specimens.

FIG. 38 Tensile test results for punched specimens from Alloy 1 beforeand after annealing demonstrating full property recovery from edgedamage by annealing. Data for EDM cut specimens for the same alloy areshown for reference.

FIG. 39 Example tensile stress-strain curves for punched specimens fromAlloy 1 with and without annealing.

FIG. 40 Tensile stress-strain curves illustrating the response of coldrolled Alloy 1 to recovery temperatures in the range between 400° C. and850° C.; a) Tensile curves, b) Yield strength.

FIG. 41 Bright-field TEM images of cold rolled ALLOY 1 samplesexhibiting the highly deformed and textured High Strength NanomodalStructure: a) Lower magnification image, b) Higher magnification image.

FIG. 42 Bright-field TEM images of ALLOY 1 samples annealed at 450° C.10 min exhibiting the highly deformed and textured High StrengthNanomodal Structure with no recrystallization occurred: a) Lowermagnification image, b) Higher magnification image.

FIG. 43 Bright-field TEM images of ALLOY 1 samples annealed at 600° C.10 min exhibiting nanoscale grains signaling the beginning ofrecrystallization: a) Lower magnification image, b) Higher magnificationimage.

FIG. 44 Bright-field TEM images of ALLOY 1 samples annealed at 650° C.10 min exhibiting larger grains indicating the higher extent ofrecrystallization: a) Lower magnification image, b) Higher magnificationimage.

FIG. 45 Bright-field TEM images of ALLOY 1 samples annealed at 700° C.10 min exhibiting recrystallized grains with a small fraction ofuntransformed area, and electron diffraction shows the recrystallizedgrains are austenite: a) Lower magnification image, b) Highermagnification image.

FIG. 46 Model Time Temperature Transformation Diagram representingresponse of the steel alloys herein to temperature at annealing. In theheating curve labeled A, recovery mechanisms are activated. In theheating curve labeled B, both recovery and recrystallization mechanismsare activated.

FIG. 47 Tensile properties of punched specimens before and afterannealing at different temperatures: a) Alloy1, b) Alloy 9, and c) Alloy12.

FIG. 48 Schematic illustration of the sample position for structuralanalysis.

FIG. 49 Alloy 1 punched E8 samples in the as-punched condition: a) Lowmagnification image showing a triangular deformation zone at the punchededge which is located on the right side of the picture. Additionallyclose up areas for the subsequent micrographs are provided, b) Highermagnification image showing the deformation zone, c) Highermagnification image showing the recrystallized structure far away fromthe deformation zone, d) Higher magnification image showing the deformedstructure in the deformation zone.

FIG. 50 Alloy 1 punched E8 samples after annealing at 650° C. for 10min: a) Low magnification image showing the deformation zone at edge,punching in upright direction. Additionally, close up areas for thesubsequent micrographs are provided: b) Higher magnification imageshowing the deformation zone, c) Higher magnification image showing therecrystallized structure far away from the deformation zone, d) Highermagnification image showing the recovered structure in the deformationzone.

FIG. 51 Alloy 1 punched E8 samples after annealing at 700° C. for 10min: a) Low magnification image showing the deformation zone at edge,punching in upright direction. Additionally, close up areas for thesubsequent micrographs are provided, b) Higher magnification imageshowing the deformation zone, c) Higher magnification image showing therecrystallized structure far away from the deformation zone, d) Highermagnification image showing the recrystallized structure in thedeformation zone.

FIG. 52 Tensile properties for specimens punched at varied speeds from:a) Alloy 1, b) Alloy 9, c) Alloy 12.

FIG. 53 HER results for Alloy 1 in a case of punched vs milled hole.

FIG. 54 Cutting plan for SEM microscopy and microhardness measurementsamples from HER tested specimens.

FIG. 55 A schematic illustration of microhardness measurement locations.

FIG. 56 Microhardness measurement profile in Alloy 1 HER tested sampleswith: a) EDM cut and b) Punched holes.

FIG. 57 Microhardness profiles for Alloy 1 in various stages ofprocessing and forming, demonstrating the progression of edge structuretransformation during hole punching and expansion.

FIG. 58 Microhardness data for HER tested samples from Alloy 1 withpunched and milled holes. Circles indicate a position of the TEM samplesin respect to hole edge.

FIG. 59 Bright field TEM image of the microstructure in the Alloy 1sheet sample before HER testing.

FIG. 60 Bright field TEM micrographs of microstructure in the HER testsample from Alloy 1 with punched hole (HER=5%) at a location of ˜1.5 mmfrom the hole edge: a) main untransformed structure; b) “pocket” ofpartially transformed structure.

FIG. 61 Bright field TEM micrographs of microstructure in the HER testsample from Alloy 1 with milled hole (HER=73.6%) at a location of ˜1.5mm from the hole edge in different areas: a) & b).

FIG. 62 Focused Ion Beam (FIB) technique used for precise sampling nearthe edge of the punched hole in the Alloy 1 sample: a) FIB techniqueshowing the general sample location of the milled TEM sample, b) Closeup view of the cut-out TEM sample with indicated location from the holeedge.

FIG. 63 Bright field TEM micrographs of microstructure in the samplefrom Alloy 1 with a punched hole at a location of ˜10 micron from thehole edge.

FIG. 64 Hole expansion ratio measurements for Alloy 1 with and withoutannealing of punched holes.

FIG. 65 Hole expansion ratio measurements for Alloy 9 with and withoutannealing of punched holes.

FIG. 66 Hole expansion ratio measurements for Alloy 12 with and withoutannealing of punched holes.

FIG. 67 Hole expansion ratio measurements for Alloy 13 with and withoutannealing of punched holes.

FIG. 68 Hole expansion ratio measurements for Alloy 17 with and withoutannealing of punched holes.

FIG. 69 Tensile performance of Alloy 1 tested with different edgeconditions. Note that tensile samples with Punched edge condition havereduced tensile performance when compared to tensile samples with wireEDM cut and punched with subsequent annealing (850° C. for 10 minutes)edge conditions.

FIG. 70 Edge formability as measured by hole expansion ratio response ofAlloy 1 as a function of edge condition. Note that holes in the Punchedcondition have lower edge formability than holes in the wire EDM cut andpunched with subsequent annealing (850° C. for 10 minutes) conditions.

FIG. 71 Punch speed dependence of Alloy 1 edge formability as a functionof punch speed, measured by hole expansion ratio. Note the consistentincrease in hole expansion ratio with increasing punch speed.

FIG. 72 Punch speed dependence of Alloy 9 edge formability as a functionof punch speed, measured by hole expansion ratio. Note the rapidincrease in hole expansion ratio up to approximately 25 mm/s punch speedfollowed by a gradual increase in hole expansion ratio.

FIG. 73 Punch speed dependence of Alloy 12 edge formability as afunction of punch speed, measured by hole expansion ratio. Note therapid increase in hole expansion ratio up to approximately 25 mm/s punchspeed followed by a continued increase in hole expansion ratio withpunch speeds of >100 mm/s.

FIG. 74 Punch speed dependence of commercial Dual Phase 980 steel edgeformability measured by hole expansion ratio. Note the hole expansionratio is consistently 21% with ±3% variance for commercial Dual Phase980 steel at all punch speeds tested.

FIG. 75 Schematic drawings of non-flat punch geometries: 6° taper(left), 7° conical (center), and conical flat (right). All dimensionsare in millimeters.

FIG. 76 Punch geometry effect on Alloy 1 at 28 mm/s, 114 mm/s, and 228mm/s punch speed. Note that for the Alloy 1, the effect of punchgeometry diminishes at 228 mm/s punch speed.

FIG. 77 Punch geometry effect on Alloy 9 at 28 mm/s, 114 mm/s, and 228mm/s punch speeds. Note that the 7° conical punch and the conical flatpunch result in the highest hole expansion ratio.

FIG. 78 Punch geometry effect on Alloy 12 at 28 mm/s, 114 mm/s, and 228mm/s punch speed. Note that the 7° conical punch results at 228 mm/spunch speed in the highest hole expansion ratio measured for all alloys.

FIG. 79 Punch geometry effect on Alloy 1 at 228 mm/s punch speed. Notethat all punch geometries result in nearly equal hole expansion ratiosof approximately 21%.

FIG. 80 Hole punch speed dependence of commercial steel grades edgeformability measured by hole expansion ratio.

FIG. 81 The post uniform elongation and hole expansion ratio correlationas predicted by [Paul S. K., J Mater Eng Perform 2014; 23:3610.] withdata for selected commercial steel grades from the same paper along withAlloy 1 and Alloy 9 data.

FIG. 82 The measured hole expansion ratio in samples from Alloy 1 as afunction of hole expansion speed.

FIG. 83 The measured hole expansion ratio in samples from Alloy 9 as afunction of hole expansion speed.

FIG. 84 The measured hole expansion ratio in samples from Alloy 12 as afunction of hole expansion speed.

FIG. 85 Images of the microstructure in the sheet from Alloy 9; a) SEMimage of the microstructure, b) Higher magnification SEM image of themicrostructure, c) Optical image of the etched surface, and d) Highermagnification optical image of the etched surface.

FIG. 86 The measured hole expansion ratio as a function of hole punchingspeed and hole expansion speed for sheet of Alloy 9.

FIG. 87 The average magnetic phases volume percent (Fe %) in the HERtested samples with different hole punching speed and hole expansionspeed as a function of the distance from the hole edge.

FIG. 88 The measured hole expansion ratio in samples from Alloy 1, Alloy9, and Alloy 12 as a function of hole preparation method.

FIG. 89 SEM images at low magnification of the cross section near thehole edge in the Alloy 1 samples with holes prepared by differentmethods prior to expansion; a) Punched hole, b) EDM cut hole, c) Milledhole, and d) Laser cut hole.

FIG. 90 SEM images at high magnification of the cross section near thehole edge in the Alloy 1 samples with holes prepared by differentmethods prior to expanding at high magnification; a) Punched hole, b)EDM cut hole, c) Milled hole, and d) Laser cut hole.

FIG. 91 SEM images at low magnification of the cross section near thehole edge in the Alloy 1 samples with holes prepared by differentmethods after expansion during HER testing; a) Punched hole, b) EDM cuthole, c) Milled hole, and d) Laser cut hole.

FIG. 92 SEM images of sample cross sections near the hole edge after HERtesting (i.e. after expansion until failure by cracking) are provided athigher magnification for samples from Alloy 1 with holes prepared bydifferent methods; a) Punched hole, b) EDM cut hole, c) Milled hole, andd) Laser cut hole.

DETAILED DESCRIPTION

Structures and Mechanisms

The steel alloys herein undergo a unique pathway of structural formationthrough specific mechanisms as illustrated in FIG. 1A and FIG. 1B.Initial structure formation begins with melting the alloy and coolingand solidifying and forming an alloy with Modal Structure (Structure #1,FIG. 1A). The Modal Structure exhibits a primarily austenitic matrix(gamma-Fe) which may contain, depending on the specific alloy chemistry,ferrite grains (alpha-Fe), martensite, and precipitates includingborides (if boron is present) and/or carbides (if carbon is present).The grain size of the Modal Structure will depend on alloy chemistry andthe solidification conditions. For example, thicker as-cast structures(e.g. thickness of greater than or equal to 2.0 mm) result in relativelyslower cooling rate (e.g. a cooling rate of less than or equal to 250K/s) and relatively larger matrix grain size. Thickness may thereforepreferably be in the range of 2.0 to 500 mm. The Modal Structurepreferably exhibits an austenitic matrix (gamma-Fe) with grain sizeand/or dendrite length from 2 to 10,000 μm and precipitates at a size of0.01 to 5.0 μm in laboratory casting. Matrix grain size and precipitatesize might be larger, up to a factor of 10 at commercial productiondepending on alloy chemistry, starting casting thickness and specificprocessing parameters. Steel alloys herein with the Modal Structure,depending on starting thickness size and the specific alloy chemistrytypically exhibits the following tensile properties, yield strength from144 to 514 MPa, ultimate tensile strength in a range from 411 to 907MPa, and total ductility from 3.7 to 24.4%.

Steel alloys herein with the Modal Structure (Structure #1, FIG. 1A) canbe homogenized and refined through the Nanophase Refinement (Mechanism#1, FIG. 1A) by exposing the steel alloy to one or more cycles of heatand stress ultimately leading to formation of the Nanomodal Structure(Structure #2, FIG. 1A). More specifically, the Modal Structure, whenformed at thickness of greater than or equal to 2.0 mm, or formed at acooling rate of less than or equal to 250 K/s, is preferably heated to atemperature of 700° C. to a temperature below the solidus temperature(T_(m)) and at strain rates of 10⁻⁶ to 10⁴ with a thickness reduction.Transformation to Structure #2 occurs in a continuous fashion throughthe intermediate Homogenized Modal Structure (Structure #1a, FIG. 1A) asthe steel alloy undergoes mechanical deformation during successiveapplication of temperature and stress and thickness reduction such aswhat can be configured to occur during hot rolling.

The Nanomodal Structure (Structure #2, FIG. 1A) has a primary austeniticmatrix (gamma-Fe) and, depending on chemistry, may additionally containferrite grains (alpha-Fe) and/or precipitates such as borides (if boronis present) and/or carbides (if carbon is present). Depending onstarting grain size, the Nanomodal Structure typically exhibits aprimary austenitic matrix (gamma-Fe) with grain size of 1.0 to 100 μmand/or precipitates at a size 1.0 to 200 nm in laboratory casting.Matrix grain size and precipitate size might be larger up to a factor of5 at commercial production depending on alloy chemistry, startingcasting thickness and specific processing parameters. Steel alloysherein with the Nanomodal Structure typically exhibit the followingtensile properties, yield strength from 264 to 574 MPa, ultimate tensilestrength in a range from 921 to 1413 MPa, and total ductility from 12.0to 77.7%. Structure #2 is preferably formed at thickness of 1 mm to 500mm.

When steel alloys herein with the Nanomodal Structure (Structure #2,FIG. 1A) are subjected to stress at ambient/near ambient temperature(e.g. 25° C. at +/−5° C.), the Dynamic Nanophase Strengthening Mechanism(Mechanism #2, FIG. 1A) is activated leading to formation of the HighStrength Nanomodal Structure (Structure #3, FIG. 1A). Preferably, thestress is at a level above the alloy's respective yield strength in arange from 250 to 600 MPa depending on alloy chemistry. The HighStrength Nanomodal structure typically exhibits a ferritic matrix(alpha-Fe) which, depending on alloy chemistry, may additionally containaustenite grains (gamma-Fe) and precipitate grains which may includeborides (if boron is present) and/or carbides (if carbon is present).Note that the strengthening transformation occurs during strain underapplied stress that defines Mechanism #2 as a dynamic process duringwhich the metastable austenitic phase (gamma-Fe) transforms into ferrite(alpha-Fe) with precipitates. Note that depending on the startingchemistry, a fraction of the austenite will be stable and will nottransform. Typically, as low as 5 volume percent and as high as 95volume percent of the matrix will transform. The High Strength NanomodalStructure typically exhibits a ferritic matrix (alpha-Fe) with matrixgrain size of 25 nm to 50 μm and precipitate grains at a size of 1.0 to200 nm in laboratory casting. Matrix grain size and precipitate sizemight be larger up to a factor of 2 at commercial production dependingon alloy chemistry, starting casting thickness and specific processingparameters. Steel alloys herein with the High Strength NanomodalStructure typically exhibits the following tensile properties, yieldstrength from 718 to 1645 MPa, ultimate tensile strength in a range from1356 to 1831 MPa, and total ductility from 1.6 to 32.8%. Structure #3 ispreferably formed at thickness of 0.2 to 25.0 mm.

The High Strength Nanomodal Structure (Structure #3, FIG. 1A and FIG.1B) has a capability to undergo Recrystallization (Mechanism #3, FIG.1B) when subjected to heating below the melting point of the alloy withtransformation of ferrite grains back into austenite leading toformation of Recrystallized Modal Structure (Structure #4, FIG. 1B).Partial dissolution of nanoscale precipitates also takes place. Presenceof borides and/or carbides is possible in the material depending onalloy chemistry. Preferred temperature ranges for a completetransformation occur from 650° C. up to the T_(m) of the specific alloy.When recrystallized, the Structure #4 contains few dislocations or twinsand stacking faults can be found in some recrystallized grains. Notethat at lower temperatures from 400 to 650° C., recovery mechanisms mayoccur. The Recrystallized Modal Structure (Structure #4, FIG. 1B)typically exhibits a primary austenitic matrix (gamma-Fe) with grainsize of 0.5 to 50 μm and precipitate grains at a size of 1.0 to 200 nmin laboratory casting. Matrix grain size and precipitate size might belarger up to a factor of 2 at commercial production depending on alloychemistry, starting casting thickness and specific processingparameters. Steel alloys herein with the Recrystallized Modal Structuretypically exhibit the following tensile properties: yield strength from197 to 1372 MPa, ultimate tensile strength in a range from 799 to 1683MPa, and total ductility from 10.6 to 86.7%.

Steel alloys herein with the Recrystallized Modal Structure (Structure#4, FIG. 1B) undergo Nanophase Refinement & Strengthening (Mechanism #4,FIG. 1B) upon stressing above yield at ambient/near ambient temperature(e.g. 25° C.+/−5° C.) that leads to formation of the Refined HighStrength Nanomodal Structure (Structure #5, FIG. 1B). Preferably thestress to initiate Mechanism #4 is at a level above yield strength in arange 197 to 1372 MPa. Similar to Mechanism #2, Nanophase Refinement &Strengthening (Mechanism #4, FIG. 1B) is a dynamic process during whichthe metastable austenitic phase transforms into ferrite with precipitateresulting generally in further grain refinement as compared to Structure#3 for the same alloy. One characteristic feature of the Refined HighStrength Nanomodal Structure (Structure #5, FIG. 1B) is that significantrefinement occurs during phase transformation in the randomlydistributed “pockets” of microstructure while other areas remainuntransformed. Note that depending on the starting chemistry, a fractionof the austenite will be stable and the area containing the stabilizedaustenite will not transform. Typically, as low as 5 volume percent andas high as 95 volume percent of the matrix in the distributed “pockets”will transform. The presence of borides (if boron is present) and/orcarbides (if carbon is present) is possible in the material depending onalloy chemistry. The untransformed part of the microstructure isrepresented by austenitic grains (gamma-Fe) with a size from 0.5 to 50μm and additionally may contain distributed precipitates with size of 1to 200 nm. These highly deformed austenitic grains contain a relativelylarge number of dislocations due to existing dislocation processesoccurring during deformation resulting in high fraction of dislocations(10⁸ to 10¹⁰ mm⁻²). The transformed part of the microstructure duringdeformation is represented by refined ferrite grains (alpha-Fe) withadditional precipitate through Nanophase Refinement & Strengthening(Mechanism #4, FIG. 1B). The size of refined grains of ferrite(alpha-Fe) varies from 50 to 2000 nm and size of precipitates is in arange from 1 to 200 nm in laboratory casting. Matrix grain size andprecipitate size might be larger up to a factor of 2 at commercialproduction depending on alloy chemistry, starting casting thickness andspecific processing parameters. The size of the “pockets” of transformedand highly refined microstructure typically varies from 0.5 to 20 μm.The volume fraction of the transformed vs untransformed areas in themicrostructure can be varied by changing the alloy chemistry includingaustenite stability from typically a 95:5 ratio to 5:95, respectively.Steel alloys herein with the Refined High Strength Nanomodal Structuretypically exhibit the following tensile properties: yield strength from718 to 1645 MPa, ultimate tensile strength in a range from 1356 to 1831MPa, and total ductility from 1.6 to 32.8%.

Steel alloys herein with the Refined High Strength Nanomodal Structure(Structure #5, FIG. 1B) may then be exposed to elevated temperaturesleading back to formation of a Recrystallized Modal Structure (Structure#4, FIG. 1B). Typical temperature ranges for a complete transformationoccur from 650° C. up to the T_(m) of the specific alloy (as illustratedin FIG. 1B) while lower temperatures from 400° C. to temperatures lessthan 650° C., activate recovery mechanisms and may cause partialrecrystallization. Stressing and heating may be repeated multiple timesto achieve desired product geometry including but not limited torelatively thin gauges of the sheet, relatively small diameter of thetube or rod, complex shape of final part, etc. with targeted properties.Final thicknesses of the material may therefore fall in the range from0.2 to 25 mm. Note that cubic precipitates may be present in the steelalloys herein at all stages with a Fm3m (#225) space group. Additionalnanoscale precipitates may be formed as a result of deformation throughDynamic Nanophase Strengthening Mechanism (Mechanism #2) and/orNanophase Refinement & Strengthening (Mechanism #4) that are representedby a dihexagonal pyramidal class hexagonal phase with a P6_(3mc) spacegroup (#186) and/or a ditrigonal dipyramidal class with a hexagonalP6bar2C space group (#190). The precipitate nature and volume fractiondepends on the alloy composition and processing history. The size ofnanoprecipitates can range from 1 nm to tens of nanometers, but in mostcases below 20 nm. Volume fraction of precipitates is generally lessthan 20%.

Mechanisms During Sheet Production Through Slab Casting

The structures and enabling mechanisms for the steel alloys herein areapplicable to commercial production using existing process flows. SeeFIG. 2. Steel slabs are commonly produced by continuous casting with amultitude of subsequent processing variations to get to the finalproduct form which is commonly coils of sheet. A detailed structuralevolution in steel alloys herein from casting to final product withrespect to each step of slab processing into sheet product isillustrated in FIG. 2.

The formation of Modal Structure (Structure #1) in steel alloys hereinoccurs during alloy solidification. The Modal Structure may bepreferably formed by heating the alloys herein at temperatures in therange of above their melting point and in a range of 1100° C. to 2000°C. and cooling below the melting temperature of the alloy, whichcorresponds to preferably cooling in the range of 1×10³ to 1×10⁻³ K/s.The as-cast thickness will be dependent on the production method withThin Slab Casting typically in the range of 20 to 150 mm in thicknessand Thick Slab Casting typically in the range of 150 to 500 mm inthickness. Accordingly, as cast thickness may fall in the range of 20 to500 mm, and at all values therein, in 1 mm increments. Accordingly, ascast thickness may be 21 mm, 22 mm, 23 mm, etc., up to 500 mm.

Hot rolling of solidified slabs from the alloys is the next processingstep with production either of transfer bars in the case of Thick SlabCasting or coils in the case of Thin Slab Casting. During this process,the Modal Structure transforms in a continuous fashion into a partialand then fully Homogenized Modal Structure (Structure #1a) throughNanophase Refinement (Mechanism #1). Once homogenization and resultingrefinement is completed, the Nanomodal Structure (Structure #2) forms.The resulting hot band coils which are a product of the hot rollingprocess is typically in the range of 1 to 20 mm in thickness.

Cold rolling is a widely used method for sheet production that isutilized to achieve targeted thickness for particular applications. ForAHSS, thinner gauges are usually targeted in the range of 0.4 to 2 mm.To achieve the finer gauge thicknesses, cold rolling can be appliedthrough multiple passes with or without intermediate annealing betweenpasses. Typical reduction per pass is 5 to 70% depending on the materialproperties and equipment capability. The number of passes before theintermediate annealing also depends on materials properties and level ofstrain hardening during cold deformation. For the steel alloys herein,the cold rolling will trigger Dynamic Nanophase Strengthening (Mechanism#2) leading to extensive strain hardening of the resultant sheet and tothe formation of the High Strength Nanomodal Structure (Structure #3).The properties of the cold rolled sheet from alloys herein will dependon the alloy chemistry and can be controlled by the cold rollingreduction to yield a fully cold rolled (i.e. hard) product or can bedone to yield a range of properties (i.e. ¼, ½, ¾ hard etc.). Dependingon the specific process flow, especially starting thickness and theamount of hot rolling gauge reduction, often annealing is needed torecover the ductility of the material to allow for additional coldrolling gauge reduction. Intermediate coils can be annealed by utilizingconventional methods such as batch annealing or continuous annealinglines. The cold deformed High Strength Nanomodal Structure (Structure#3) for the steel alloys herein will undergo Recrystallization(Mechanism #3) during annealing leading to the formation of theRecrystallized Modal Structure (Structure #4). At this stage, therecrystallized coils can be a final product with advanced propertycombination depending on the alloy chemistry and targeted markets. In acase when even thinner gauges of the sheet are required, recrystallizedcoils can be subjected to further cold rolling to achieve targetedthickness that can be realized by one or multiple cycles of coldrolling/annealing. Additional cold deformation of the sheet from alloysherein with Recrystallized Modal Structure (Structure #4) leads tostructural transformation into Refined High Strength Nanomodal Structure(Structure #5) through Nanophase Refinement and Strengthening (Mechanism#4). As a result, fully hard coils with final gauge and Refined HighStrength Nanomodal Structure (Structure #5) can be formed or, in thecase of annealing as a last step in the cycle, coils of the sheet withfinal gauge and Recrystallized Modal Structure (Structure #4) can alsobe produced. When coils of recrystallized sheet from alloys hereinutilized for finished part production by any type of cold deformationsuch as cold stamping, hydroforming, roll forming etc., Refined HighStrength Nanomodal Structure (Structure #5) will be present in the finalproduct/parts. The final products may be in many different formsincluding sheet, plate, strips, pipes, and tubes and a myriad of complexparts made through various metalworking processes.

Mechanisms for Edge Formability

The cyclic nature of these phase transformations going fromRecrystallized Modal Structure (Structure #4) to Refined High StrengthNanomodal Structure (Structure #5) and then back to Recrystallized ModalStructure (Structure #4) is one of the unique phenomenon and features ofsteel alloys herein. As described earlier, this cyclic feature isapplicable during commercial manufacturing of the sheet, especially forAHSS where thinner gauge thicknesses are required (e.g. thickness in therange of 0.2 to 25 mm). Furthermore, these reversibility mechanisms areapplicable for the widespread industrial usage of the steel alloysherein. While exhibiting exceptional combinations of bulk sheetformability as is demonstrated by the tensile and bend properties inthis application for the steel alloys herein, the unique cycle featureof the phase transformations is enabling for edge formability, which canbe a significant limiting factor for other AHSS. Table 1 below providesa summary of the structure and performance features through stressingand heating cycles available through Nanophase Refinement andStrengthening (Mechanism #4). How these structures and mechanisms can beharnessed to produce exceptional combinations of both bulk sheet andedge formability will be subsequently described herein.

TABLE 1 Structures and Performance Through Stressing/Heating CyclesMechanism Structure #5 Structure #4 Refined High Strength NanomodalStructure Property Recrystallized Modal Structure UntransformedTransformed “pockets” Structure Recrystallization Retained austeniticNanophase Refinement & Formation occurring at elevated grainsStrengthening mechanism temperatures in cold occurring through workedmaterial application of mechanical stress in distributed micro-structural “pockets” Transformations Recrystallization of coldPrecipitation Stress induced austenite deformed iron matrix optionaltransformation into ferrite and precipitates Enabling Phases Austenite,optionally Austenite, optionally Ferrite, optionally ferrite,precipitates precipitates austenite, precipitates Matrix Grain Size 0.5to 50 μm   0.5 to 50 μm 50 to 2000 nm Precipitate Size 1 to 200 nm   1to 200 nm 1 to 200 nm Tensile Response Actual with properties achievedActual with properties achieved based on formation of the structurebased on formation of the structure and fraction of transformation andfraction of transformation Yield Strength 197 to 1372 MPa  718 to 1645MPa Ultimate Tensile Strength 799 to 1683 MPa 1356 to 1831 MPa TotalElongation 6.6 to 86.7%   1.6 to 32.8% 

Main Body

The chemical composition of the alloys herein is shown in Table 2 whichprovides the preferred atomic ratios utilized.

TABLE 2 Alloy Chemical Composition Alloy Fe Cr Ni Mn Cu B Si C Alloy 175.75 2.63 1.19 13.86 0.65 0.00 5.13 0.79 Alloy 2 73.99 2.63 1.19 13.181.55 1.54 5.13 0.79 Alloy 3 77.03 2.63 3.79 9.98 0.65 0.00 5.13 0.79Alloy 4 78.03 2.63 5.79 6.98 0.65 0.00 5.13 0.79 Alloy 5 79.03 2.63 7.793.98 0.65 0.00 5.13 0.79 Alloy 6 78.53 2.63 3.79 8.48 0.65 0.00 5.130.79 Alloy 7 79.53 2.63 5.79 5.48 0.65 0.00 5.13 0.79 Alloy 8 80.53 2.637.79 2.48 0.65 0.00 5.13 0.79 Alloy 9 74.75 2.63 1.19 14.86 0.65 0.005.13 0.79 Alloy 10 75.25 2.63 1.69 13.86 0.65 0.00 5.13 0.79 Alloy 1174.25 2.63 1.69 14.86 0.65 0.00 5.13 0.79 Alloy 12 73.75 2.63 1.19 15.860.65 0.00 5.13 0.79 Alloy 13 77.75 2.63 1.19 11.86 0.65 0.00 5.13 0.79Alloy 14 74.75 2.63 2.19 13.86 0.65 0.00 5.13 0.79 Alloy 15 73.75 2.633.19 13.86 0.65 0.00 5.13 0.79 Alloy 16 74.11 2.63 2.19 13.86 1.29 0.005.13 0.79 Alloy 17 72.11 2.63 2.19 15.86 1.29 0.00 5.13 0.79 Alloy 1878.25 2.63 0.69 11.86 0.65 0.00 5.13 0.79 Alloy 19 74.25 2.63 1.19 14.861.15 0.00 5.13 0.79 Alloy 20 74.82 2.63 1.50 14.17 0.96 0.00 5.13 0.79Alloy 21 75.75 1.63 1.19 14.86 0.65 0.00 5.13 0.79 Alloy 22 77.75 2.631.19 13.86 0.65 0.00 3.13 0.79 Alloy 23 76.54 2.63 1.19 13.86 0.65 0.005.13 0.00 Alloy 24 67.36 10.70 1.25 10.56 1.00 5.00 4.13 0.00 Alloy 2571.92 5.45 2.10 8.92 1.50 6.09 4.02 0.00 Alloy 26 61.30 18.90 6.80 0.900.00 5.50 6.60 0.00 Alloy 27 71.62 4.95 4.10 6.55 2.00 3.76 7.02 0.00Alloy 28 62.88 16.00 3.19 11.36 0.65 0.00 5.13 0.79 Alloy 29 72.50 2.630.00 15.86 1.55 1.54 5.13 0.79 Alloy 30 80.19 0.00 0.95 13.28 1.66 2.250.88 0.79 Alloy 31 77.65 0.67 0.08 13.09 1.09 0.97 2.73 3.72 Alloy 3278.54 2.63 1.19 13.86 0.65 0.00 3.13 0.00 Alloy 33 83.14 1.63 8.68 0.001.00 4.76 0.00 0.79 Alloy 34 75.30 2.63 1.34 14.01 0.80 0.00 5.13 0.79Alloy 35 74.85 2.63 1.49 14.16 0.95 0.00 5.13 0.79

As can be seen from the above, the alloys herein are iron based metalalloys, having greater than or equal to 50 at. % Fe. More preferably,the alloys herein can be described as comprising, consisting essentiallyof, or consisting of the following elements at the indicated atomicpercent: Fe (61.30 to 83.14 at. %); Si (0 to 7.02 at. %); Mn (0 to 15.86at. %); B (0 to 6.09 at. %); Cr (0 to 18.90 at. %); Ni (0 to 8.68 at.%); Cu (0 to 2.00 at. %); C (0 to 3.72 at. %). In addition, it can beappreciated that the alloys herein are such that they comprise Fe and atleast four or more, or five or more, or six or more elements selectedfrom Si, Mn, B, Cr, Ni, Cu or C. Most preferably, the alloys herein aresuch that they comprise, consist essentially of, or consist of Fe at alevel of 50 at. % or greater along with Si, Mn, B, Cr, Ni, Cu and C.

Alloy Laboratory Processing

Laboratory processing of the alloys in Table 2 was done to model eachstep of industrial production but on a much smaller scale. Key steps inthis process include the following: casting, tunnel furnace heating, hotrolling, cold rolling, and annealing.

Casting

Alloys were weighed out into charges ranging from 3,000 to 3,400 gramsusing commercially available ferroadditive powders with known chemistryand impurity content according to the atomic ratios in Table 2. Chargeswere loaded into a zirconia coated silica crucibles which was placedinto an Indutherm VTC800V vacuum tilt casting machine. The machine thenevacuated the casting and melting chambers and backfilled with argon toatmospheric pressure several times prior to casting to prevent oxidationof the melt. The melt was heated with a 14 kHz RF induction coil untilfully molten, approximately 5.25 to 6.5 minutes depending on the alloycomposition and charge mass. After the last solids were observed to meltit was allowed to heat for an additional 30 to 45 seconds to providesuperheat and ensure melt homogeneity. The casting machine thenevacuated the melting and casting chambers, tilted the crucible andpoured the melt into a 50 mm thick, 75 to 80 mm wide, and 125 mm deepchannel in a water cooled copper die. The melt was allowed to cool undervacuum for 200 seconds before the chamber was filled with argon toatmospheric pressure. Example pictures of laboratory cast slabs from twodifferent alloys are shown in FIG. 3.

Tunnel Furnace Heating

Prior to hot rolling, laboratory slabs were loaded into a LuciferEHS3GT-B18 furnace to heat. The furnace set point varies between 1100°C. to 1250° C. depending on alloy melting point. The slabs were allowedto soak for 40 minutes prior to hot rolling to ensure they reach thetarget temperature. Between hot rolling passes the slabs are returned tothe furnace for 4 minutes to allow the slabs to reheat.

Hot Rolling

Pre-heated slabs were pushed out of the tunnel furnace into a Fenn Model061 2 high rolling mill. The 50 mm slabs were preferably hot rolled for5 to 8 passes though the mill before being allowed to air cool. Afterthe initial passes each slab had been reduced between 80 to 85% to afinal thickness of between 7.5 and 10 mm. After cooling each resultantsheet was sectioned and the bottom 190 mm was hot rolled for anadditional 3 to 4 passes through the mill, further reducing the platebetween 72 to 84% to a final thickness of between 1.6 and 2.1 mm.Example pictures of laboratory cast slabs from two different alloysafter hot rolling are shown in FIG. 4.

Cold Rolling

After hot rolling resultant sheets were media blasted with aluminumoxide to remove the mill scale and were then cold rolled on a Fenn Model061 2 high rolling mill. Cold rolling takes multiple passes to reducethe thickness of the sheet to a targeted thickness of typically 1.2 mm.Hot rolled sheets were fed into the mill at steadily decreasing rollgaps until the minimum gap is reached. If the material has not yet hitthe gauge target, additional passes at the minimum gap were used until1.2 mm thickness was achieved. A large number of passes were applied dueto limitations of laboratory mill capability. Example pictures of coldrolled sheets from two different alloys are shown in FIG. 5.

Annealing

After cold rolling, tensile specimens were cut from the cold rolledsheet via wire electrical discharge machining (EDM). These specimenswere then annealed with different parameters listed in Table 3.Annealing 1a, 1b, 2b were conducted in a Lucifer 7HT-K12 box furnace.Annealing 2a and 3 was conducted in a Camco Model G-ATM-12FL furnace.Specimens which were air normalized were removed from the furnace at theend of the cycle and allowed to cool to room temperature in air. For thefurnace cooled specimens, at the end of the annealing the furnace wasshut off to allow the sample to cool with the furnace. Note that theheat treatments were selected for demonstration but were not intended tobe limiting in scope. High temperature treatments up to just below themelting points for each alloy are possible.

TABLE 3 Annealing Parameters An- Temper- nealing Heating ature DwellCooling Atmosphere 1a Preheated 850° C.  5 min Air Normalized Air +Argon Furnace 1b Preheated 850° C.  10 min Air Normalized Air + ArgonFurnace 2a 20° C./hr 850° C. 360 min 45° C./hr to Hydrogen + 500° C.then Argon Furnace Cool 2b 20° C./hr 850° C. 360 min 45° C./hr to Air +Argon 500° C. then Air Normalized 3 20° C./hr 1200° C.  120 min FurnaceCool Hydrogen + ArgonAlloy Properties

Thermal analysis of the alloys herein was performed on as-solidifiedcast slabs using a Netzsch Pegasus 404 Differential Scanning calorimeter(DSC). Samples of alloys were loaded into alumina crucibles which werethen loaded into the DSC. The DSC then evacuated the chamber andbackfilled with argon to atmospheric pressure. A constant purge of argonwas then started, and a zirconium getter was installed in the gas flowpath to further reduce the amount of oxygen in the system. The sampleswere heated until completely molten, cooled until completely solidified,then reheated at 10° C./min through melting. Measurements of thesolidus, liquidus, and peak temperatures were taken from the secondmelting in order to ensure a representative measurement of the materialin an equilibrium state. In the alloys listed in Table 2, melting occursin one or multiple stages with initial melting from ˜1111° C. dependingon alloy chemistry and final melting temperature up to ˜1476° C. (Table4). Variations in melting behavior reflect complex phase formation atsolidification of the alloys depending on their chemistry.

TABLE 4 Differential Thermal Analysis Data for Melting Behavior SolidusLiquidus Melting Melting Melting Temperature Temperature Peak #1 Peak #2Peak #3 Alloy (° C.) (° C.) (° C.) (° C.) (° C.) Alloy 1 1390 1448 1439Alloy 2 1157 1410 1177 1401 Alloy 3 1411 1454 1451 Alloy 4 1400 14601455 Alloy 5 1415 1467 1464 Alloy 6 1416 1462 1458 Alloy 7 1421 14671464 Alloy 8 1417 1469 1467 Alloy 9 1385 1446 1441 Alloy 10 1383 14421437 Alloy 11 1384 1445 1442 Alloy 12 1385 1443 1435 Alloy 13 1401 14591451 Alloy 14 1385 1445 1442 Alloy 15 1386 1448 1441 Alloy 16 1384 14391435 Alloy 17 1376 1442 1435 Alloy 18 1395 1456 1431 1449 1453 Alloy 191385 1437 1432 Alloy 20 1374 1439 1436 Alloy 21 1391 1442 1438 Alloy 221408 1461 1458 Alloy 23 1403 1452 1434 1448 Alloy 24 1219 1349 1246 13141336 Alloy 25 1186 1335 1212 1319 Alloy 26 1246 1327 1268 1317 Alloy 271179 1355 1202 1344 Alloy 28 1158 1402 1176 1396 Alloy 29 1159 1448 11681439 Alloy 30 1111 1403 1120 1397 Alloy 31 1436 1475 1464 Alloy 32 14361476 1464 Alloy 33 1153 1418 1178 1411 Alloy 34 1397 1448 1445 Alloy 351394 1444 1441

The density of the alloys was measured on 9 mm thick sections of hotrolled material using the Archimedes method in a specially constructedbalance allowing weighing in both air and distilled water. The densityof each alloy is tabulated in Table 5 and was found to be in the rangefrom 7.57 to 7.89 g/cm³. The accuracy of this technique is ±0.01 g/cm³.

TABLE 5 Density of Alloys Density Alloy (g/cm³) Alloy 1 7.78 Alloy 27.74 Alloy 3 7.82 Alloy 4 7.84 Alloy 5 7.76 Alloy 6 7.83 Alloy 7 7.79Alloy 8 7.71 Alloy 9 7.77 Alloy 10 7.78 Alloy 11 7.77 Alloy 12 7.77Alloy 13 7.80 Alloy 14 7.78 Alloy 15 7.79 Alloy 16 7.79 Alloy 17 7.77Alloy 18 7.79 Alloy 19 7.77 Alloy 20 7.78 Alloy 21 7.78 Alloy 22 7.87Alloy 23 7.81 Alloy 24 7.67 Alloy 25 7.71 Alloy 26 7.57 Alloy 27 7.67Alloy 28 7.73 Alloy 29 7.89 Alloy 30 7.78 Alloy 31 7.89 Alloy 32 7.89Alloy 33 7.78 Alloy 34 7.77 Alloy 35 7.78

Tensile properties were measured on an Instron 3369 mechanical testingframe using Instron's Bluehill control software. All tests wereconducted at room temperature, with the bottom grip fixed and the topgrip set to travel upwards at a rate of 0.012 mm/s. Strain data wascollected using Instron's Advanced Video Extensometer. Tensileproperties of the alloys listed in Table 2 after annealing withparameters listed in Table 3 are shown below in Table 6 to Table 10. Theultimate tensile strength values may vary from 799 to 1683 MPa withtensile elongation from 6.6 to 86.7%. The yield strength is in a rangefrom 197 to 978 MPa. The mechanical characteristic values in the steelalloys herein will depend on alloy chemistry and processing conditions.The variation in heat treatment additionally illustrates the propertyvariations possible through processing a particular alloy chemistry.

TABLE 6 Tensile Data for Selected Alloys after Heat Treatment 1a YieldUltimate Tensile Strength Tensile Strength Elongation Alloy (MPa) (MPa)(%) Alloy 1 443 1212 51.1 458 1231 57.9 422 1200 51.9 Alloy 2 484 127848.3 485 1264 45.5 479 1261 48.7 Alloy 3 458 1359 43.9 428 1358 43.7 4621373 44.0 Alloy 4 367 1389 36.4 374 1403 39.1 364 1396 32.1 Alloy 5 5101550 16.5 786 1547 18.1 555 1552 16.2 Alloy 6 418 1486 34.3 419 147535.2 430 1490 37.3 Alloy 7 468 1548 20.2 481 1567 20.3 482 1545 19.3Alloy 8 851 1664 13.6 848 1683 14.0 859 1652 12.9 Alloy 9 490 1184 58.0496 1166 59.1 493 1144 56.6 Alloy 10 472 1216 60.5 481 1242 58.7 4701203 55.9 Alloy 11 496 1158 65.7 498 1155 58.2 509 1154 68.3 Alloy 12504 1084 48.3 515 1105 70.8 518 1106 66.9 Alloy 13 478 1440 41.4 4861441 40.7 455 1424 42.0 Alloy 22 455 1239 48.1 466 1227 55.4 460 123757.9 Alloy 23 419 1019 48.4 434 1071 48.7 439 1084 47.5 Alloy 28 583 93261.5 594 937 60.8 577 930 61.0 Alloy 29 481 1116 60.0 481 1132 55.4 4861122 56.8 Alloy 30 349 1271 42.7 346 1240 36.2 340 1246 42.6 Alloy 31467 1003 36.0 473 996 29.9 459 988 29.5 Alloy 32 402 1087 44.2 409 106146.1 420 1101 44.1

TABLE 7 Tensile Data for Selected Alloys after Heat Treatment 1b YieldUltimate Tensile Strength Tensile Strength Elongation Alloy (MPa) (MPa)(%) Alloy 1 487 1239 57.5 466 1269 52.5 488 1260 55.8 Alloy 2 438 123249.7 431 1228 49.8 431 1231 49.4 Alloy 9 522 1172 62.6 466 1170 61.9 4621168 61.3 Alloy 12 471 1115 63.3 458 1102 69.3 454 1118 69.1 Alloy 13452 1408 40.5 435 1416 42.5 432 1396 46.0 Alloy 14 448 1132 64.4 4431151 60.7 436 1180 54.3 Alloy 15 444 1077 66.9 438 1072 65.3 423 107570.5 Alloy 16 433 1084 67.5 432 1072 66.8 423 1071 67.8 Alloy 17 420 94674.6 421 939 77.0 425 961 74.9 Alloy 19 496 1124 67.4 434 1118 64.8 4351117 67.4 Alloy 20 434 1154 58.3 457 1188 54.9 448 1187 60.5 Alloy 21421 1201 54.3 427 1185 59.9 431 1191 47.8 Alloy 24 554 1151 23.5 5381142 24.3 562 1151 24.3 Alloy 25 500 1274 16.0 502 1271 15.8 483 128016.3 Alloy 26 697 1215 20.6 723 1187 21.3 719 1197 21.5 Alloy 27 5381385 20.6 574 1397 20.9 544 1388 21.8 Alloy 33 978 1592 6.6 896 1596 7.2953 1619 7.5 Alloy 34 467 1227 56.7 476 1232 52.7 462 1217 51.6 Alloy 35439 1166 56.3 438 1166 59.0 440 1177 58.3

TABLE 8 Tensile Data for Selected Alloys after Heat Treatment 2a YieldUltimate Tensile Strength Tensile Strength Elongation Alloy (MPa) (MPa)(%) Alloy 2 367 1174 46.2 369 1193 45.1 367 1179 50.2 Alloy 30 391 111855.7 389 1116 60.5 401 1113 59.5 Alloy 32 413 878 17.6 399 925 20.5 384962 21.0 Alloy 31 301 1133 37.4 281 1125 38.7 287 1122 39.0

TABLE 9 Tensile Data for Selected Alloys after Heat Treatment 2b YieldUltimate Tensile Strength Tensile Strength Elongation Alloy (MPa) (MPa)(%) Alloy 1 396 1093 31.2 383 1070 30.4 393 1145 34.7 Alloy 2 378 123349.4 381 1227 48.3 366 1242 47.7 Alloy 3 388 1371 41.3 389 1388 42.6Alloy 4 335 1338 21.7 342 1432 30.1 342 1150 17.3 Alloy 5 568 1593 15.2595 1596 13.1 735 1605 14.6 Alloy 6 399 1283 17.5 355 1483 24.8 386 147123.8 Alloy 7 605 1622 16.3 639 1586 15.2 Alloy 8 595 1585 13.6 743 162314.1 791 1554 13.9 Alloy 9 381 1125 53.3 430 1111 44.8 369 1144 51.1Alloy 10 362 1104 37.8 369 1156 43.5 Alloy 11 397 1103 52.4 390 108650.9 402 1115 50.4 Alloy 12 358 1055 64.7 360 1067 64.4 354 1060 62.9Alloy 13 362 982 17.3 368 961 16.3 370 989 17.0 Alloy 14 385 1165 59.0396 1156 55.5 437 1155 57.9 Alloy 15 357 1056 70.3 354 1046 68.2 3581060 70.7 Alloy 16 375 1094 67.6 384 1080 63.4 326 1054 65.2 Alloy 17368 960 77.2 370 955 77.9 358 951 75.9 Alloy 18 326 1136 17.3 338 119219.1 327 1202 18.5 Alloy 19 386 1134 64.5 378 1100 60.5 438 1093 52.5Alloy 20 386 1172 56.2 392 1129 42.0 397 1186 57.8 Alloy 21 363 114149.0 Alloy 22 335 1191 45.7 322 1189 41.5 348 1168 34.5 Alloy 23 3981077 44.3 367 1068 44.8 Alloy 24 476 1149 28.0 482 1154 25.9 495 114526.2 Alloy 25 452 1299 16.0 454 1287 15.8 441 1278 15.1 Alloy 26 6191196 26.6 615 1189 26.2 647 1193 26.1 Alloy 27 459 1417 17.3 461 141016.8 457 1410 17.1 Alloy 28 507 879 52.3 498 874 42.5 493 880 44.7 Alloy32 256 1035 42.3 257 1004 42.1 257 1049 34.8 Alloy 33 830 1494 8.4 8621521 8.1 877 1519 8.8 Alloy 34 388 1178 59.8 384 1197 57.7 370 1177 59.1Alloy 35 367 1167 58.5 369 1167 58.4 375 1161 59.7

TABLE 10 Tensile Data for Selected Alloys after Heat Treatment 3 YieldUltimate Tensile Strength Tensile Strength Elongation Alloy (MPa) (MPa)(%) Alloy 1 238 1142 47.6 233 1117 46.3 239 1145 53.0 Alloy 3 266 133838.5 N/A 1301 37.7 N/A 1291 35.6 Alloy 4 N/A 1353 27.7 N/A 1337 26.1 N/A1369 29.0 Alloy 5 511 1462 12.5 558 1399 10.6 Alloy 6 311 1465 24.6 3081467 21.8 308 1460 25.0 Alloy 7 727 1502 12.5 639 1474 11.3 685 152012.4 Alloy 8 700 1384 12.3 750 1431 13.3 Alloy 9 234 1087 55.0 240 107056.4 242 1049 58.3 Alloy 10 229 1073 50.6 228 1082 56.5 229 1077 54.2Alloy 11 232 1038 63.8 232 1009 62.4 228 999 66.1 Alloy 12 229 979 65.6228 992 57.5 222 963 66.2 Alloy 13 277 1338 37.3 261 1352 35.9 272 135334.9 Alloy 14 228 1074 58.5 239 1077 54.1 230 1068 49.1 Alloy 15 206 99160.9 208 1024 58.9 Alloy 16 199 1006 57.7 242 987 53.4 208 995 57.0Alloy 17 222 844 72.6 197 867 64.9 213 869 66.5 Alloy 18 288 1415 32.6300 1415 32.1 297 1421 29.6 Alloy 19 225 1032 58.5 213 1019 61.1 2141017 58.4 Alloy 20 233 1111 57.3 227 1071 53.0 230 1091 49.4 Alloy 21238 1073 50.6 228 1069 56.5 246 1110 52.0 Alloy 22 217 1157 47.0 2361154 46.8 218 1154 47.7 Alloy 23 208 979 45.4 204 984 43.4 204 972 38.9Alloy 28 277 811 86.7 279 802 86.0 277 799 82.0 Alloy 32 203 958 33.3206 966 39.5 210 979 36.3 Alloy 34 216 1109 52.8 230 1144 55.9 231 112352.3 Alloy 35 230 1104 51.7 231 1087 59.0 220 1084 54.4

CASE EXAMPLES Case Example #1: Structural Development Pathway in Alloy 1

A laboratory slab with thickness of 50 mm was cast from Alloy 1 that wasthen laboratory processed by hot rolling, cold rolling and annealing at850° C. for 5 min as described in Main Body section of currentapplication. Microstructure of the alloy was examined at each step ofprocessing by SEM, TEM and x-ray analysis.

For SEM study, the cross section of the slab samples was ground on SiCabrasive papers with reduced grit size, and then polished progressivelywith diamond media paste down to 1 μm. The final polishing was done with0.02 μm grit SiO₂ solution. Microstructures were examined by SEM usingan EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMTInc. To prepare TEM specimens, the samples were first cut by EDM, andthen thinned by grinding with pads of reduced grit size every time.Further thinning to make foils of 60 to 70 μm thickness was done bypolishing with 9 μm, 3 μm and 1 μm diamond suspension solutionrespectively. Discs of 3 mm in diameter were punched from the foils andthe final polishing was completed with electropolishing using a twin-jetpolisher. The chemical solution used was a 30% nitric acid mixed inmethanol base. In case of insufficient thin area for TEM observation,the TEM specimens may be ion-milled using a Gatan Precision IonPolishing System (PIPS). The ion-milling usually is done at 4.5 keV, andthe inclination angle is reduced from 4° to 2° to open up the thin area.The TEM studies were done using a JEOL 2100 high-resolution microscopeoperated at 200 kV. X-ray diffraction was done using a PANalyticalX'Pert MPD diffractometer with a Cu Kα x-ray tube and operated at 45 kVwith a filament current of 40 mA. Scans were run with a step size of0.01° and from 25° to 95° two-theta with silicon incorporated to adjustfor instrument zero angle shift. The resulting scans were thensubsequently analyzed using Rietveld analysis using Siroquant software.

Modal Structure was formed in the Alloy 1 slab with 50 mm thicknessafter solidification. The Modal Structure (Structure #1) is representedby a dendritic structure that is composed of several phases. In FIG. 6a, the backscattered SEM image shows the dendritic arms that are shown indark contrast while the matrix phase is in bright contrast. Note thatsmall casting pores are found as exhibited (black holes) in the SEMmicrograph. TEM studies show that the matrix phase is primarilyaustenite (gamma-Fe) with stacking faults (FIG. 6b ). The presence ofstacking faults indicates a face-centered-cubic structure (austenite).TEM also suggests that other phases could be formed in the ModalStructure. As shown in FIG. 6c , a dark phase is found that identifiedas a ferrite phase with body-centered cubic structure (alpha-Fe)according to selected electron diffraction pattern. X-ray diffractionanalysis shows that the Modal Structure of the Alloy 1 containsaustenite, ferrite, iron manganese compound and some martensite (FIG.7). Generally, austenite is the dominant phase in the Alloy 1 ModalStructure, but other factors such as the cooling rate during commercialproduction may influence the formation of secondary phases such asmartensite with varying volume fraction.

TABLE 11 X-ray Diffraction Data for Alloy 1 After Solidification (ModalStructure) Phases Identified Phase Details γ - Fe Structure: Cubic Spacegroup #: 225 (Fm3m) LP: a = 3.583 Å α - Fe Structure: Cubic Space group#: 229 (Im3m) LP: a = 2.876 Å Martensite Structure: Tetragonal Spacegroup #: 139 (I4/mmm) LP: a = 2.898 Å c = 3.018 Å Iron manganesecompound Structure: Cubic Space group #: 225 (Fm3m) LP: a = 4.093 Å

Deformation of the Alloy 1 with the Modal Structure (Structure #1, FIG.1A) at elevated temperature induces homogenization and refinement ofModal Structure. Hot rolling was applied in this case but otherprocesses including but not limited to hot pressing, hot forging, hotextrusion can achieve the similar effect. During hot rolling, thedendrites in the Modal Structure are broken up and refined, leadinginitially to the Homogenized Modal Structure (Structure #1a, FIG. 1A)formation. The refinement during the hot rolling occurs through theNanophase Refinement (Mechanism #1, FIG. 1A) along with dynamicrecrystallization. The Homogenized Modal Structure can be progressivelyrefined by applying the hot rolling repetitively, leading to theNanomodal Structure (Structure #2, FIG. 1A) formation. FIG. 8a shows thebackscattered SEM micrograph of Alloy 1 after being hot rolled from 50mm to ˜1.7 mm at 1250° C. It can be seen that blocks of tens of micronsin size are resulted from the dynamic recrystallization during the hotrolling, and the interior of the grains is relatively smooth indicatingless amount of defects. TEM further reveals that sub-grains of less thanseveral hundred nanometers in size are formed, as shown in FIG. 8b .X-ray diffraction analysis shows that the Nanomodal Structure of theAlloy 1 after hot rolling contains mainly austenite, with other phasessuch as ferrite and the iron manganese compound as shown in FIG. 9 andTable 12.

TABLE 12 X-ray Diffraction Data for Alloy 1 After Hot Rolling (NanomodalStructure) Phases Identified Phase Details γ - Fe Structure: Cubic Spacegroup #: 225 (Fm3m) LP: a = 3.595 Å α - Fe Structure: Cubic Space group#: 229 (Im3m) LP: a = 2.896 Å Iron manganese compound Structure: CubicSpace group #: 225 (Fm3m) LP: a = 4.113 Å

Further deformation at ambient temperature (i.e., cold deformation) ofthe Alloy 1 with the Nanomodal Structure causes transformation into HighStrength Nanomodal Structure (Structure #3, FIG. 1A) through the DynamicNanophase Strengthening (Mechanism #2, FIG. 1A). The cold deformationcan be achieved by cold rolling and, tensile deformation, or other typeof deformation such as punching, extrusion, stamping, etc. During thecold deformation, depending on alloy chemistries, a large portion ofaustenite in the Nanomodal Structure is transformed to ferrite withgrain refinement. FIG. 10a shows the backscattered SEM micrograph ofcold rolled Alloy 1. Compared to the smooth grains in the NanomodalStructure after hot rolling, the cold deformed grains are roughindicating severe plastic deformation within the grains. Depending onalloy chemistry, deformation twins can be produced in some alloysespecially by cold rolling, as displayed in FIG. 10a . FIG. 10b showsthe TEM micrograph of the microstructure in cold rolled Alloy 1. It canbe seen that in addition to dislocations generated by the deformation,refined grains due to phase transformation can also be found. The bandedstructure is related to the deformation twins caused by the coldrolling, corresponding to these in FIG. 10a . X-ray diffraction showsthat the High Strength Nanomodal Structure of the Alloy 1 after coldrolling contains a significant amount of ferrite phase in addition tothe retained austenite and the iron manganese compound as shown in FIG.11 and Table.

TABLE 13 X-ray Diffraction Data for Alloy 1 after Cold Rolling (HighStrength Nanomodal Structure) Phases Identified Phase Details γ - FeStructure: Cubic Space group #: 225 (Fm3m) LP: a = 3.588 Å α - FeStructure: Cubic Space group #: 229 (Im3m) LP: a = 2.871 Å Ironmanganese compound Structure: Cubic Space group #: 225 (Fm3m) LP: a =4.102 Å

Recrystallization occurs upon heat treatment of the cold deformed Alloy1 with High Strength Nanomodal Structure (Structure #3, FIGS. 1A and 1B)that transforms into Recrystallized Modal Structure (Structure #4, FIG.1B). The TEM images of the Alloy 1 after annealing are shown in, FIG.12. As it can be seen, equiaxed grains with sharp and straightboundaries are present in the structure and the grains are free ofdislocations, which is characteristic feature of recrystallization.Depending on the annealing temperature, the size of recrystallizedgrains can range from 0.5 to 50 μm. In addition, as shown in electrondiffraction shows that austenite is the dominant phase afterrecrystallization. Annealing twins are occasionally found in the grains,but stacking faults are most often seen. The formation of stackingfaults shown in the TEM image is typical for face-centered-cubic crystalstructure of austenite. Backscattered SEM micrographs in FIG. 13 showthe equiaxed recrystallized grains with the size of less than 10 μm,consistent with TEM. The different contrast of grains (dark or bright)seen on SEM images suggests that the crystal orientation of the grainsis random, since the contrast in this case is mainly originated from thegrain orientation. As a result, any texture formed by the previous colddeformation is eliminated. X-ray diffraction shows that theRecrystallized Modal Structure of the Alloy 1 after annealing containsprimarily austenite phase, with a small amount of ferrite and the ironmanganese compound as shown in FIG. 14 and Table 14.

TABLE 14 X-ray Diffraction Data for Alloy 1 After Annealing(Recrystallized Modal Structure) Phases Identified Phase Details γ - FeStructure: Cubic Space group #: 225 (Fm3m) LP: a = 3.597 Å α - FeStructure: Cubic Space group #: 229 (Im3m) LP: a = 2.884 Å Ironmanganese compound Structure: Cubic Space group #: 225 (Fm3m) LP: a =4.103 Å

When the Alloy 1 with Recrystallized Modal Structure (Structure #4, FIG.1B) is subjected to deformation at ambient temperature, NanophaseRefinement & Strengthening (Mechanism #4, FIG. 1B) is activated leadingto formation of the Refined High Strength Nanomodal Structure (Structure#5, FIG. 1B). In this case, deformation was a result of tensile testingand gage section of the tensile sample after testing was analyzed. FIG.15 shows the bright-field TEM micrographs of the microstructure in thedeformed Alloy 1. Compared to the matrix grains that were initiallyalmost dislocation-free in the Recrystallized Modal Structure afterannealing, the application of stress generates a high density ofdislocations within the matrix grains. At the end of tensile deformation(with a tensile elongation greater than 50%), accumulation of largenumber of dislocations is observed in the matrix grains. As shown inFIG. 15a , in some areas (for example the area at the lower part of theFIG. 15a ), dislocations form a cell structure and the matrix remainsaustenitic. In other areas, where the dislocation density issufficiently high, transformation is induced from austenite to ferrite(for example the upper and right part of the FIG. 15a ) that results insubstantial structure refinement. FIG. 15b shows local “pocket” of thetransformed refined microstructure and selected area electrondiffraction pattern corresponds to ferrite. Structural transformationinto Refined High Strength Nanomodal Structure (Structure #5, FIG. 1B)in the randomly distributed “pockets” is a characteristic feature of thesteel alloys herein. FIG. 16 shows the backscattered SEM images of theRefined High Strength Nanomodal Structure. Compared to theRecrystallized Modal Structure, the boundaries of matrix grains becomeless apparent, and the matrix is obviously deformed. Although thedetails of deformed grains cannot be revealed by SEM, the change causedby the deformation is enormous compared to the Recrystallized ModalStructure that was demonstrated in TEM images. X-ray diffraction showsthat the Refined High Strength Nanomodal Structure of the Alloy 1 aftertensile deformation contains a significant amount of ferrite andaustenite phases. Very broad peaks of ferrite phase (alpha-Fe) are seenin the XRD pattern, suggesting significant refinement of the phase. Theiron manganese compound is also present. Additionally, a hexagonal phasewith space group #186 (P6_(3mc)) was identified in the gage section ofthe tensile sample as shown in FIG. 17 and Table 15.

TABLE 15 X-ray Diffraction Data for Alloy 1 After Tensile Deformation(Refined High Strength Nanomodal Structure) Phases Identified PhaseDetails γ - Fe Structure: Cubic Space group #: 225 (Fm3m) LP: a = 3.586Å α - Fe Structure: Cubic Space group #: 229 (Im3m) LP: a = 2.873 Å Ironmanganese compound Structure: Cubic Space group #: 225 (Fm3m) LP: a =4.159 Å Hexagonal phase 1 Structure: Hexagonal Space group #: 186(P6₃mc) LP: a = 3.013 Å, c = 6.183 Å

This Case Example demonstrates that alloys listed in Table 2 includingAlloy 1 exhibit a structural development pathway with novel enablingmechanisms illustrated in FIGS. 1A and 1B leading to uniquemicrostructures with nanoscale features.

Case Example #2 Structural Development Pathway in Alloy 2

Laboratory slab with thickness of 50 mm was cast from Alloy 2 that wasthen laboratory processed by hot rolling, cold rolling and annealing at850° C. for 10 min as described in Main Body section of currentapplication. Microstructure of the alloy was examined at each step ofprocessing by SEM, TEM and x-ray analysis.

For SEM study, the cross section of the slab samples was ground on SiCabrasive papers with reduced grit size, and then polished progressivelywith diamond media paste down to 1 μm. The final polishing was done with0.02 μm grit SiO₂ solution. Microstructures were examined by SEM usingan EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMTInc. To prepare TEM specimens, the samples were first cut with EDM, andthen thinned by grinding with pads of reduced grit size every time.Further thinning to make foils to ˜60 μm thickness was done by polishingwith 9 μm, 3 μm and 1 μm diamond suspension solution respectively. Discsof 3 mm in diameter were punched from the foils and the final polishingwas fulfilled with electropolishing using a twin-jet polisher. Thechemical solution used was a 30% nitric acid mixed in methanol base. Incase of insufficient thin area for TEM observation, the TEM specimensmay be ion-milled using a Gatan Precision Ion Polishing System (PIPS).The ion-milling usually is done at 4.5 keV, and the inclination angle isreduced from 4° to 2° to open up the thin area. The TEM studies weredone using a JEOL 2100 high-resolution microscope operated at 200 kV.X-ray diffraction was done using a Panalytical X'Pert MPD diffractometerwith a Cu Kα x-ray tube and operated at 45 kV with a filament current of40 mA. Scans were run with a step size of 0.01° and from 25° to 95°two-theta with silicon incorporated to adjust for instrument zero angleshift. The resulting scans were then subsequently analyzed usingRietveld analysis using Siroquant software.

Modal Structure (Structure #1, FIG. 1A) is formed in Alloy 2 slab castat 50 mm thick, which is characterized by dendritic structure. Due tothe presence of a boride phase (M₂B), the dendritic structure is moreevident than in Alloy 1 where borides are absent. FIG. 18a shows thebackscattered SEM of Modal Structure that exhibits a dendritic matrix(in bright contrast) with borides at the boundary (in dark contrast).TEM studies show that the matrix phase is composed of austenite(gamma-Fe) with stacking faults (FIG. 18b ). Similar to Alloy 1, thepresence of stacking faults indicates the matrix phase is austenite.Also shown in TEM is the boride phase that appears dark in. FIG. 18b atthe boundary of austenite matrix phase. X-ray diffraction analysis datain. FIG. 19 and Table 16 shows that the Modal Structure containsaustenite, M₂B, ferrite, and iron manganese compound. Similar to Alloy1, austenite is the dominant phase in the Alloy 2 Modal Structure, butother phases may be present depending on alloy chemistry.

TABLE 16 X-ray Diffraction Data for Alloy 2 After Solidification (ModalStructure) Phases Identified Phase Details γ - Fe Structure: Cubic Spacegroup #: 225 (Fm3m) LP: a = 3.577 Å α - Fe Structure: Cubic Space group#: 229 (Im3m) LP: a = 2.850 Å M₂B Structure: Tetragonal Space group #:140 (I4/mcm) LP: a = 5.115 Å, c = 4.226 Å Iron manganese compoundStructure: Cubic Space group #: 225 (Fm3m) LP: a = 4.116 Å

Following the flowchart in FIG. 1A, deformation of the Alloy 2 with theModal Structure (Structure #1, FIG. 1A) at elevated temperature induceshomogenization and refinement of Modal Structure. Hot rolling wasapplied in this case but other processes including but not limited tohot pressing, hot forging, hot extrusion can achieve a similar effect.During the hot rolling, the dendrites in the Modal Structure are brokenup and refined, leading initially to the Homogenized Modal Structure(Structure #1a, FIG. 1A) formation. The refinement during the hotrolling occurs through the Nanophase Refinement (Mechanism #1, FIG. 1A)along with dynamic recrystallization. The Homogenized Modal Structurecan be progressively refined by applying the hot rolling repetitively,leading to the Nanomodal Structure (Structure #2, FIG. 1A) formation.FIG. 20a shows the backscattered SEM micrograph of hot rolled Alloy 2.Similar to Alloy 1, the dendritic Modal Structure is homogenized whilethe boride phase is randomly distributed in the matrix. TEM shows thatthe matrix phase is partially recrystallized as a result of dynamicrecrystallization during hot rolling, as shown in FIG. 20b . The matrixgrains are on the order of 500 nm, which is finer than in Alloy 1 due tothe pinning effect of borides. X-ray diffraction analysis shows that theNanomodal Structure of Alloy 2 after hot rolling contains mainlyaustenite phase and M₂B, with other phases such as ferrite and ironmanganese compound as shown in FIG. 21 and Table 17.

TABLE 17 X-ray Diffraction Data for Alloy 2 After Hot Rolling (NanomodalStructure) Phases Identified Phase Details γ - Fe Structure: Cubic Spacegroup #: 225 (Fm3m) LP: a = 3.598 Å α - Fe Structure: Cubic Space group#: 229 (Im3m) LP: a = 2.853 Å M₂B Structure: Tetragonal Space group #:140 (I4/mcm) LP: a = 5.123 Å, c = 4.182 Å Iron manganese compoundStructure: Cubic Space group #: 225 (Fm3m) LP: a = 4.180 Å

Deformation of the Alloy 2 with the Nanomodal Structure but at ambienttemperature (i.e., cold deformation) leads to formation of High StrengthNanomodal Structure (Structure #3, FIG. 1A) through the DynamicNanophase Strengthening (Mechanism #2, FIG. 1A). The cold deformationcan be achieved by cold rolling, tensile deformation, or other type ofdeformation such as punching, extrusion, stamping, etc. Similarly inAlloy 2 during cold deformation, a great portion of austenite in theNanomodal Structure is transformed to ferrite with grain refinement.FIG. 22a shows the backscattered SEM micrograph of the microstructure inthe cold rolled Alloy 2. Deformation is concentrated in the matrix phasearound the boride phase. FIG. 22b shows the TEM micrograph of the coldrolled Alloy 2. Refined grains can be found due to the phasetransformation. Although deformation twins are less evident in SEMimage, TEM shows that they are generated after the cold rolling, similarto Alloy 1. X-ray diffraction shows that the High Strength NanomodalStructure of the Alloy 2 after cold rolling contains a significantamount of ferrite phase in addition to the M₂B, retained austenite and anew hexagonal phase with space group #186 (P6_(3mc)) as shown in FIG. 23and Table 18.

TABLE 18 X-ray Diffraction Data for Alloy 2 After Cold Rolling (HighStrength Nanomodal Structure) Phases Identified Phase Details γ - FeStructure: Cubic Space group #: 225 (Fm3m) LP: a = 3.551 Å α - FeStructure: Cubic Space group #: 229 (Im3m) LP: a = 2.874 Å M₂BStructure: Tetragonal Space group #: 140 (I4/mcm) LP: a = 5.125 Å, c =4.203 Å Hexagonal phase Structure: Hexagonal Space group #: 186 (P6₃mc)LP: a = 2.962 Å, c = 6.272 Å

Recrystallization occurs upon annealing of the cold deformed Alloy 2with High Strength Nanomodal Structure (Structure #3, FIGS. 1A and 1B)that transforms into Recrystallized Modal Structure (Structure #4, FIG.1B). The recrystallized microstructure of the Alloy 2 after annealing isshown by TEM images in FIG. 24. As it can be seen, equiaxed grains withsharp and straight boundaries are present in the structure and thegrains are free of dislocations, which is a characteristic feature ofrecrystallization. The size of recrystallized grains is generally lessthan 5 μm due to the pinning effect of boride phase, but larger grainsare possible at higher annealing temperatures. Moreover, electrondiffraction shows that austenite is the dominant phase afterrecrystallization and stacking faults are present in the austenite, asshown in FIG. 24b . The formation of stacking faults also indicatesformation of face-centered-cubic austenite phase. Backscattered SEMmicrographs in FIG. 25 show the equiaxed recrystallized grains with thesize of less than 5 μm, with boride phase randomly distributed. Thedifferent contrast of grains (dark or bright) seen on SEM imagessuggests that the crystal orientation of the grains is random, since thecontrast in this case is mainly originated from the grain orientation.As a result, any texture formed by the previous cold deformation iseliminated. X-ray diffraction shows that the Recrystallized ModalStructure of the Alloy 2 after annealing contains primarily austenitephase, with M₂B, a small amount of ferrite, and a hexagonal phase withspace group #186 (P6_(3mc)) as shown in FIG. 26 and Table 19.

TABLE 19 X-ray Diffraction Data for Alloy 2 After Annealing(Recrystallized Modal Structure) Phases Identified Phase Details γ - FeStructure: Cubic Space group #: 225 (Fm3m) LP: a = 3.597 Å α - FeStructure: Cubic Space group #: 229 (Im3m) LP: a = 2.878 Å M₂BStructure: Tetragonal Space group #: 140 (I4/mcm) LP: a = 5.153 Å, c =4.170 Å Hexagonal phase Structure: Hexagonal Space group #: 186(P6_(3mc)) LP: a = 2.965 Å, c = 6.270 Å

Deformation of Recrystallized Modal Structure (Structure #4, FIG. 1B)leads to formation of the Refined High Strength Nanomodal Structure(Structure #5, FIG. 1B) through Nanophase Refinement & Strengthening(Mechanism #4, FIG. 1B). In this case, deformation was a result oftensile testing and the gage section of the tensile sample after testingwas analyzed. FIG. 27 shows the micrographs of microstructure in thedeformed Alloy 2. Similar to Alloy 1, the initially dislocation-freematrix grains in the Recrystallized Modal Structure after annealing arefilled with a high density of dislocations upon the application ofstress, and the accumulation of dislocations in some grains activatesthe phase transformation from austenite to ferrite, leading tosubstantial refinement. As shown in FIG. 27a , refined grains of 100 to300 nm in size are shown in a local “pocket” where transformationoccurred from austenite to ferrite. Structural transformation intoRefined High Strength Nanomodal Structure (Structure #5, FIG. 1B) in the“pockets” of matrix grains is a characteristic feature of the steelalloys herein. FIG. 27b shows the backscattered SEM images of theRefined High Strength Nanomodal Structure. Similarly, the boundaries ofmatrix grains become less apparent after the matrix is deformed. X-raydiffraction shows that a significant amount of austenite transformed toferrite although the four phases remain as in the Recrystallized ModalStructure. The transformation resulted in formation of Refined HighStrength Nanomodal Structure of the Alloy 2 after tensile deformation.Very broad peaks of ferrite phase (α-Fe) are seen in the XRD pattern,suggesting significant refinement of the phase. As in Alloy 1, a newhexagonal phase with space group #186 (P6_(3mc)) was identified in thegage section of the tensile sample as shown in FIG. 28 and Table 20.

TABLE 20 X-ray Diffraction Data for Alloy 2 After Tensile Deformation(Refined High Strength Nanomodal Structure) Phases Identified PhaseDetails γ - Fe Structure: Cubic Space group #: 225 (Fm3m) LP: a = 3.597Å α - Fe Structure: Cubic Space group #: 229 (Im3m) LP: a = 2.898 Å M₂BStructure: Tetragonal Space group #: 140 (I4/mcm) LP: a = 5.149 Å, c =4.181 Å Hexagonal phase Structure: Hexagonal Space group #: 186(P6_(3mc)) LP: a = 2.961 Å, c = 6.271 Å

This Case Example demonstrates that alloys listed in Table 2 includingAlloy 2 exhibit a structural development pathway with the mechanismsillustrated in FIGS. 1A and 1B leading to unique microstructures withnanoscale features.

Case Example #3 Tensile Properties at Each Step of Processing

Slabs with thickness of 50 mm were laboratory cast from the alloyslisted in Table 21 according to the atomic ratios provided in Table 2and laboratory processed by hot rolling, cold rolling and annealing at850° C. for 10 min as described in Main Body section of currentapplication. Tensile properties were measured at each step of processingon an Instron 3369 mechanical testing frame using Instron's Bluehillcontrol software. All tests were conducted at room temperature, with thebottom grip fixed and the top grip set to travel upwards at a rate of0.012 mm/s. Strain data was collected using Instron's Advanced VideoExtensometer.

Alloys were weighed out into charges ranging from 3,000 to 3,400 gramsusing commercially available ferroadditive powders with known chemistryand impurity content according to the atomic ratios in Table 2. Chargeswere loaded into zirconia coated silica crucibles which were placed intoan Indutherm VTC800V vacuum tilt casting machine. The machine thenevacuated the casting and melting chambers and backfilled with argon toatmospheric pressure several times prior to casting to prevent oxidationof the melt. The melt was heated with a 14 kHz RF induction coil untilfully molten, approximately 5.25 to 6.5 minutes depending on the alloycomposition and charge mass. After the last solids were observed to meltit was allowed to heat for an additional 30 to 45 seconds to providesuperheat and ensure melt homogeneity. The casting machine thenevacuated the melting and casting chambers and tilted the crucible andpoured the melt into a 50 mm thick, 75 to 80 mm wide, and 125 mm deepchannel in a water cooled copper die. The melt was allowed to cool undervacuum for 200 seconds before the chamber was filled with argon toatmospheric pressure. Tensile specimens were cut from as-cast slabs bywire EDM and tested in tension. Results of tensile testing are shown inTable 21. As it can be seen, ultimate tensile strength of the alloysherein in as-cast condition varies from 411 to 907 MPa. The tensileelongation varies from 3.7 to 24.4%. Yield strength is measured in arange from 144 to 514 MPa.

Prior to hot rolling, laboratory cast slabs were loaded into a LuciferEHS3GT-B18 furnace to heat. The furnace set point varies between 1000°C. to 1250° C. depending on alloy melting point. The slabs were allowedto soak for 40 minutes prior to hot rolling to ensure they reach thetarget temperature. Between hot rolling passes the slabs are returned tothe furnace for 4 minutes to allow the slabs to reheat. Pre-heated slabswere pushed out of the tunnel furnace into a Fenn Model 061 2 highrolling mill. The 50 mm casts are hot rolled for 5 to 8 passes throughthe mill before being allowed to air cool defined as first campaign ofhot rolling. After this campaign the slab thickness was reduced between80.4 to 87.4%. After cooling, the resultant sheet samples were sectionedto 190 mm in length. These sections were hot rolled for an additional 3passes through the mill with reduction between 73.1 to 79.9% to a finalthickness of between 2.1 and 1.6 mm. Detailed information on hot rollingconditions for each alloy herein is provided in Table 22. Tensilespecimens were cut from hot rolled sheets by wire EDM and tested intension. Results of tensile testing are shown in Table 22. After hotrolling, ultimate tensile strength of the alloys herein varies from 921to 1413 MPa. The tensile elongation varies from 12.0 to 77.7%. Yieldstrength is measured in a range from 264 to 574 MPa. See, Structure 2 inFIG. 1A.

After hot rolling, resultant sheets were media blasted with aluminumoxide to remove the mill scale and were then cold rolled on a Fenn Model061 2 high rolling mill. Cold rolling takes multiple passes to reducethe thickness of the sheet to targeted thickness, generally 1.2 mm. Hotrolled sheets were fed into the mill at steadily decreasing roll gapsuntil the minimum gap is reached. If the material has not yet hit thegauge target, additional passes at the minimum gap were used until thetargeted thickness was reached. Cold rolling conditions with the numberof passes for each alloy herein are listed in Table 23. Tensilespecimens were cut from cold rolled sheets by wire EDM and tested intension. Results of tensile testing are shown in Table 23. Cold rollingleads to significant strengthening with ultimate tensile strength in therange from 1356 to 1831 MPa. The tensile elongation of the alloys hereinin cold rolled state varies from 1.6 to 32.1%. Yield strength ismeasured in a range from 793 to 1645 MPa. It is anticipated that higherultimate tensile strength and yield strength can be achieved in alloysherein by larger cold rolling reduction (>40%) that in our case islimited by laboratory mill capability. With more rolling force, it isanticipated that ultimate tensile strength could be increased to atleast 2000 MPa and yield strength to at least 1800 MPa.

Tensile specimens were cut from cold rolled sheet samples by wire EDMand annealed at 850° C. for 10 min in a Lucifer 7HT-K12 box furnace.Samples were removed from the furnace at the end of the cycle andallowed to cool to room temperature in air. Results of tensile testingare shown in Table 24. As it can be seen, recrystallization duringannealing of the alloys herein results in property combinations withultimate tensile strength in the range from 939 to 1424 MPa and tensileelongation from 15.8 to 77.0%. Yield strength is measured in a rangefrom 420 to 574 MPa.

FIG. 29 to FIG. 31 represent plotted data at each processing step forAlloy 1, Alloy 13, and Alloy 17, respectively.

TABLE 21 Tensile Properties of Alloys in As-Cast State Yield UltimateTensile Strength Tensile Strength Elongation Alloy (MPa) (MPa) (%) Alloy1 289 527 10.4 288 548 9.3 260 494 8.4 Alloy 2 244 539 10.4 251 592 11.6249 602 13.1 Alloy 13 144 459 4.6 156 411 4.5 163 471 5.7 Alloy 17 223562 24.4 234 554 20.7 235 585 23.3 Alloy 24 396 765 8.3 362 662 5.7 404704 7.0 Alloy 25 282 668 5.1 329 753 5.0 288 731 5.5 Alloy 25 471 7884.1 514 907 6.0 483 815 3.7 Alloy 27 277 771 3.7 278 900 4.9 267 798 4.5Alloy 34 152 572 11.1 168 519 11.6 187 545 12.9 Alloy 35 164 566 15.9172 618 16.6 162 569 16.4

TABLE 22 Tensile Properties of Alloys in Hot Rolled State First SecondCam- Cam- Ultimate Tensile paign paign Yield Tensile Elon- Reduc- Reduc-Strength Strength gation Alloy Condition tion tion (MPa) (MPa) (%) Alloy1 Hot 80.5%, 75.1%, 273 1217 50.0 Rolled 6 Passes 3 Passes 264 1216 52.195.2% 285 1238 52.7 Alloy 2 Hot 87.4%, 73.1%, 480 1236 45.3 Rolled 7Passes 3 Passes 454 1277 41.9 96.6% 459 1219 48.2 Alloy 13 Hot 81.1%,79.8%, 287 1116 18.8 Rolled 6 Passes 3 Passes 274 921 15.3 96.0% 2931081 19.3 Alloy 17 Hot 81.2%, 79.1%, 392 947 73.3 Rolled 6 Passes 3Passes 363 949 74.8 96.1% 383 944 77.7 Alloy 24 Hot 81.1%, 79.9%, 5191176 21.4 Rolled, 6 Passes 3 Passes 521 1088 18.2 96.2% 508 1086 17.9Alloy 25 Hot 81.0%, 79.4%, 502 1105 12.4 Rolled 6 Passes 3 Passes 5241100 12.3 96.1% 574 1077 12.0 Alloy 27 Hot 80.4%, 78.9%, 508 1401 20.9Rolled, 6 Passes 3 Passes 534 1405 22.4 95.9% 529 1413 19.7 Alloy 34 Hot80.7%, 80.1%, 346 1188 56.5 Rolled, 6 Passes 3 Passes 323 1248 58.796.2% 303 1230 53.4 Alloy 35 Hot 80.8%, 79.9%, 327 1178 63.3 Rolled, 6Passes 3 Passes 317 1170 61.2 96.1% 305 1215 59.6

TABLE 23 Tensile Properties of Alloys in Cold Rolled State YieldUltimate Tensile Strength Tensile Strength Elongation Alloy Condition(MPa) (MPa) (%) Alloy 1 Cold Rolled 20.3%, 798 1492 28.5 4 Passes 7931482 32.1 Cold Rolled 39.6%, 1109 1712 21.4 29 Passes 1142 1726 23.01203 1729 21.2 Alloy 2 Cold Rolled 28.5%, 966 1613 13.4 5 Passes 9981615 15.4 1053 1611 20.6 Cold Rolled 39.1%, 1122 1735 20.3 19 passes1270 1744 18.3 Alloy 13 Cold Rolled 36.0%, 1511 1824 9.5 24 Passes 14241803 7.7 1361 1763 5.1 Alloy 17 Cold Rolled 38.5%, 1020 1357 24.2 8Passes 1007 1356 24.9 1071 1357 24.9 Alloy 24 Cold Rolled 38.2%, 13631584 1.9 23 Passes 1295 1601 2.5 1299 1599 3.0 Alloy 25 Cold Rolled38.0%, 1619 1761 1.9 42 Passes 1634 1741 1.7 1540 1749 1.6 Alloy 27 ColdRolled 39.4%, 1632 1802 2.7 40 Passes 1431 1804 4.1 1645 1831 4.1 Alloy34 Cold Rolled 35.%, 1099 1640 14.7 14 Passes 840 1636 17.5 1021 166118.5 Alloy 35 Cold Rolled 35.5%, 996 1617 23.8 12 Passes 1012 1614 24.51020 1616 23.3

TABLE 24 Tensile Properties of Alloys in Annealed State Yield UltimateTensile Strength Tensile Strength Elongation Alloy (MPa) (MPa) (%) Alloy1 436 1221 54.9 443 1217 56.0 431 1216 59.7 Alloy 2 438 1232 49.7 4311228 49.8 431 1231 49.4 484 1278 48.3 485 1264 45.5 479 1261 48.7 Alloy13 441 1424 41.7 440 1412 41.4 429 1417 42.7 Alloy 17 420 946 74.6 421939 77.0 425 961 74.9 Alloy 24 554 1151 23.5 538 1142 24.3 562 1151 24.3Alloy 25 500 1274 16.0 502 1271 15.8 483 1280 16.3 Alloy 27 538 138520.6 574 1397 20.9 544 1388 21.8 Alloy 27 467 1227 56.7 476 1232 52.7462 1217 51.6 Alloy 27 439 1166 56.3 438 1166 59.0 440 1177 58.3

This Case Example demonstrates that due to the unique mechanisms andstructural pathway shown in FIGS. 1A and 1B, the structures andresulting properties in steel alloys herein can vary widely leading tothe development of 3^(rd) Generation AHSS.

Case Example #4 Cyclic Reversibility During Cold Rolling andRecrystallization

Slabs with thickness of 50 mm were laboratory cast from Alloy 1 andAlloy 2 according to the atomic ratios provided in Table 2 and hotrolled into sheets with final thickness of 2.31 mm for Alloy 1 sheet and2.35 mm for Alloy 2 sheet. Casting and hot rolling procedures aredescribed in Main Body section of current application. Resultant hotrolled sheet from each alloy was used for demonstration of cyclicstructure/property reversibility through cold rolling/annealing cycles.

Hot rolled sheet from each alloy was subjected to three cycles of coldrolling and annealing. Sheet thicknesses before and after hot rollingand cold rolling reduction at each cycle are listed in Table 25.Annealing at 850° C. for 10 min was applied after each cold rolling.Tensile specimens were cut from the sheet in the initial hot rolledstate and at each step of the cycling. Tensile properties were measuredon an Instron 3369 mechanical testing frame using Instron's Bluehillcontrol software. All tests were conducted at room temperature, with thebottom grip fixed and the top grip set to travel upwards at a rate of0.012 mm/s. Strain data was collected using Instron's Advanced VideoExtensometer.

The results of tensile testing are plotted in FIG. 32 for Alloy 1 andAlloy 2 showing that cold rolling results in significant strengtheningof both alloys at each cycle with average ultimate tensile strength of1500 MPa in Alloy 1 and 1580 MPa in Alloy 2. Both cold rolled alloysshow a loss in ductility as compared to the hot rolled state. However,annealing after cold rolling at each cycle results in tensile propertyrecovery to the same level with high ductility.

Tensile properties for each tested sample are listed in Table 26 andTable 27 for Alloy 1 and Alloy 2, respectively. As it can be seen, Alloy1 has ultimate tensile strength from 1216 to 1238 MPa in hot rolledstate with ductility from 50.0 to 52.7% and yield strength from 264 to285 MPa. In cold rolled state, the ultimate tensile strength wasmeasured in the range from 1482 to 1517 MPa at each cycle. Ductility wasfound consistently in the range from 28.5 to 32.8% with significantlyhigher yield strength of 718 to 830 MPa as compared to that in hotrolled condition. Annealing at each cycle resulted in restoration of theductility to the range from 47.7 to 59.7% with ultimate tensile strengthfrom 1216 to 1270 MPa. Yield strength after cold rolling and annealingis lower than that after cold rolling and was measured in the range from431 to 515 MPa that is however higher than that in initial hot rolledcondition.

Similar results with property reversibility between cold rolled andannealed material through cycling were observed for Alloy 2 (FIG. 32b ).In initial hot rolled state, Alloy 2 has ultimate tensile strength from1219 to 1277 MPa with ductility from 41.9 to 48.2% and yield strengthfrom 454 to 480 MPa. Cold rolling at each cycle results in the materialstrengthening to the ultimate tensile strength from 1553 to 1598 MPawith ductility reduction to the range from 20.3 to 24.1%. Yield strengthwas measured from 912 to 1126 MPa. After annealing at each cycle, Alloy2 has ultimate tensile strength from 1231 to 1281 MPa with ductilityfrom 46.9 to 53.5%. Yield strength in Alloy 2 after cold rolling andannealing at each cycle is similar to that in hot rolled condition andvaries from 454 to 521 MPa.

TABLE 25 Sample Thickness and Cycle Reduction at Cold Rolling StepsInitial Final Cycle Rolling Thickness Thickness Reduction Alloy Cycle(mm) (mm) (%) Alloy 1 1 2.35 1.74 26.0 2 1.74 1.32 24.1 3 1.32 1.02 22.7Alloy 2 1 2.31 1.85 19.9 2 1.85 1.51 18.4 3 1.51 1.22 19.2

TABLE 26 Tensile Properties of Alloy 1 Through Cold Rolling/AnnealingCycles 1st Cycle 2nd Cycle 3rd Cycle Hot Cold Cold Cold Property RolledRolled Annealed Rolled Annealed Rolled Annealed Ultimate 1217 1492 12211497 1239 1517 1270 Tensile 1216 1482 1217 1507 1269 1507 1262 Strength1238 * 1216 1503 1260 1507 1253 (MPa) Yield 273 798 436 775 487 820 508Strength 264 793 443 718 466 796 501 (MPa) 285 * 431 830 488 809 515Tensile 50.0 28.5 54.9 32.8 57.5 32.1 50.5 Elongation 52.1 32.1 56.029.4 52.5 30.2 47.7 (%) 52.7 * 59.7 30.9 55.8 30.5 55.5 * Specimensslipped in the grips/data is not available

TABLE 27 Tensile Properties of Alloy 2 Through Cold Rolling/AnnealingCycles 1st Cycle 2nd Cycle 3rd Cycle Hot Cold Cold Cold Property RolledRolled Annealed Rolled Annealed Rolled Annealed Ultimate 1236 1579 12501553 1243 1596 1231 Tensile 1277 * 1270 1568 1255 1589 1281 Strength1219 * 1240 1566 1242 1598 1269 (MPa) Yield 480 1126 466 983 481 1006475 Strength 454 * 468 969 521 978 507 (Mpa) 459 * 454 912 497 1011 518Tensile 45.3    20.3 53.0 24.1 51.1 22.3 46.9 Elongation 41.9 * 51.223.1 52.3 23.2 53.5 (%) 48.2 * 51.1 21.6 49.9 21.0 47.9 * Specimensslipped in the grips/data is not available

This Case Example demonstrates that the High Strength NanomodalStructure (Structure #3, FIG. 1A) that forms in the alloys listed inTable 2 after cold rolling can be recrystallized by applying an annealto produce a Recrystallized Modal Structure (Structure #4, FIG. 1B).This structure can be further deformed through cold rolling or othercold deformation approaches to undergo Nanophase Refinement andStrengthening (Mechanism #4, FIG. 1B) leading to formation of theRefined High Strength Nanomodal Structure (Structure #5, FIG. 1B). TheRefined High Strength Nanomodal Structure (Structure #5, FIG. 1B) can inturn be recrystallized and the process can be started over with fullstructure/property reversibility through multiple cycles. The abilityfor the mechanisms to be reversible enables the production of finergauges which are important for weight reduction when using AHSS as wellas property recovery after any damage caused by deformation.

Case Example #5 Bending Ability

Slabs with thickness of 50 mm were laboratory cast from selected alloyslisted in Table 28 according to the atomic ratios provided in Table 2and laboratory processed by hot rolling, cold rolling and annealing at850° C. for 10 min as described in Main Body section of currentapplication. Resultant sheet from each alloy with final thickness of˜1.2 mm and Recrystallized Modal Structure (Structure #4, FIG. 1B) wasused to evaluate bending response of alloys herein.

Bend tests were performed using an Instron 5984 tensile test platformwith an Instron W-6810 guided bend test fixture according tospecifications outlined in the ISO 7438 International Standard Metallicmaterials—Bend test (International Organization for Standardization,2005). Test specimens were cut by wire EDM to a dimension of 20 mm×55mm×sheet thickness. No special edge preparation was done to the samples.Bend tests were performed using an Instron 5984 tensile test platformwith an Instron W-6810 guided bend test fixture. Bend tests wereperformed according to specifications outlined in the ISO 7438International Standard Metallic materials—Bend test (InternationalOrganization for Standardization, 2005).

The test was performed by placing the test specimen on the fixturesupports and pushing with a former as shown in FIG. 33.

The distance between supports, l, was fixed according to ISO 7438 duringthe test at:

$\begin{matrix}{l = {\left( {D + {3a}} \right) \pm \frac{a}{2}}} & {{Equation}\mspace{14mu} 1}\end{matrix}$

Prior to bending, the specimens were lubricated on both sides with 3 in1 oil to reduce friction with the test fixture. This test was performedwith a 1 mm diameter former. The former was pushed downward in themiddle of the supports to different angles up to 180° or until a crackappeared. The bending force was applied slowly to permit free plasticflow of the material. The displacement rate was calculated based on thespan gap of each test in order to have a constant angular rate andapplied accordingly.

Absence of cracks visible without the use of magnifying aids wasconsidered evidence that the test piece withstood the bend test. If acrack was detected, the bend angle was measured manually with a digitalprotractor at the bottom of the bend. The test specimen was then removedfrom the fixture and examined for cracking on the outside of the bendradius. The onset of cracking could not be conclusively determined fromthe force-displacement curves and was instead easily determined bydirect observation with illumination from a flashlight.

Results of the bending response of the alloys herein are listed in Table28 including initial sheet thickness, former radius to sheet thicknessratio (r/t) and maximum bend angle before cracking. All alloys listed inthe Table 28 did not show cracks at 90° bend angle. The majority of thealloys herein have capability to be bent at 180° angle without cracking.Example of the samples from Alloy 1 after bend testing to 180° is shownin FIG. 34.

TABLE 7 Bend Test Results for Selected Alloys Former Maximum DiameterThickness Bend Angle Alloy (mm) (mm) r/t (°) Alloy 1 0.95 1.185 0.401180 1.200 0.396 180 1.213 0.392 180 1.223 0.388 180 1.181 0.402 1801.187 0.400 180 1.189 0.399 180 1.206 0.394 180 Alloy 2 0.95 1.225 0.388180 1.230 0.386 180 1.215 0.391 180 1.215 0.391 180 1.215 0.391 1801.224 0.388 180 1.208 0.393 180 1.208 0.393 180 Alloy 3 0.95 1.212 0.392180 1.186 0.401 180 1.201 0.396 180 Alloy 4 0.95 1.227 0.387 180 1.1850.401 180 1.187 0.400 180 Alloy 5 0.95 1.199 0.396 110 1.196 0.397 90Alloy 6 0.95 1.259 0.377 160 1.202 0.395 165 1.206 0.394 142 Alloy 70.95 1.237 0.384 104 1.236 0.384 90 Alloy 9 0.95 1.278 0.372 180 1.1970.397 180 1.191 0.399 180 Alloy 10 0.95 1.226 0.387 180 1.208 0.393 1001.208 0.393 180 1.205 0.394 180 Alloy 11 0.95 1.240 0.383 180 1.2140.391 180 1.205 0.394 180 Alloy 12 0.95 1.244 0.382 180 1.215 0.391 1801.205 0.394 180 Alloy 13 0.95 1.222 0.389 180 1.191 0.399 180 1.1880.400 180 Alloy 14 0.95 1.239 0.383 180 1.220 0.389 180 1.214 0.391 180Alloy 15 0.95 1.247 0.381 180 1.224 0.388 180 1.224 0.388 180 Alloy 160.95 1.244 0.382 180 1.224 0.388 180 1.199 0.396 180 Alloy 17 0.95 1.2330.385 180 1.213 0.392 180 1.203 0.395 180 Alloy 18 0.95 1.222 0.389 1601.218 0.390 135 Alloy 19 0.95 1.266 0.375 180 1.243 0.382 180 1.2420.382 180 Alloy 20 0.95 1.242 0.382 180 1.222 0.389 180 1.220 0.389 180Alloy 21 0.95 1.255 0.378 180 1.228 0.387 180 1.229 0.386 180 Alloy 220.95 1.240 0.383 180 1.190 0.399 180 1.190 0.399 180 Alloy 23 0.95 1.1900.399 180 1.199 0.396 180 1.193 0.398 180 Alloy 28 0.95 1.222 0.389 1801.206 0.394 180 1.204 0.395 180 Alloy 29 0.95 1.219 0.390 180 1.2170.390 180 1.206 0.394 180 Alloy 30 0.95 1.215 0.391 180 1.212 0.392 1751.200 0.396 180 Alloy 31 0.95 1.211 0.392 150 1.209 0.393 131 Alloy 320.95 1.222 0.389 180 1.221 0.389 180 1.210 0.393 180

In order to be made into complex parts for automobile and other uses, anAHSS needs to exhibit both bulk sheet formability and edge sheetformability. This Case Example demonstrates good bulk sheet formabilityof the alloys in Table 2 through bend testing.

Case Example #6 Punched Edge vs EDM Cut Tensile Properties

Slabs with thickness of 50 mm were laboratory cast from selected alloyslisted in Table 2 according to the atomic ratios provided in Table 2 andlaboratory processed by hot rolling, cold rolling and annealing at 850°C. for 10 min as described herein. Resultant sheet from each alloy withfinal thickness of 1.2 mm and Recrystallized Modal Structure (Structure#4, FIG. 1B) were used to evaluate the effect of edge damage on alloyproperties by cutting tensile specimens by wire electrical dischargemachining (wire-EDM) (which represents the control situation or relativelack of shearing and formation of an edge without a compromise inmechanical properties) and by punching (to identify a mechanicalproperty loss due to shearing). It should be appreciated that shearing(imposition of a stress coplanar with a material cross-section) mayoccur herein by a number of processing options, such as piercing,perforating, cutting or cropping (cutting off of an end of a given metalpart).

Tensile specimens in the ASTM E8 geometry were prepared using both wireEDM cutting and punching. Tensile properties were measured on an Instron5984 mechanical testing frame using Instron's Bluehill control software.All tests were conducted at room temperature, with the bottom grip fixedand the top grip set to travel upwards at a rate of 0.012 mm/s. Straindata was collected using Instron's Advanced Video Extensometer. Tensiledata is shown in Table 29 and illustrated in FIG. 35a for selectedalloys. Decrease in properties is observed for all alloys tested but thelevel of this decrease varies significantly depending on alloychemistry. Table 30 summarizes a comparison of ductility in punchedsamples as compared to that in the wire EDM cut samples. In FIG. 35bcorresponding tensile curves are shown for the selected alloydemonstrating mechanical behavior as a function of austenite stability.For selected alloys herein, austenite stability is highest in Alloy 12that shows high ductility and lowest in Alloy 13 that shows highstrength. Correspondingly, Alloy 12 demonstrated lowest loss inductility in punched specimens vs EDM cut (29.7% vs 60.5%, Table 30)while Alloy 13 demonstrated highest loss in ductility in punchedspecimens vs EDM cut (5.2% vs 39.1%, Table 30). High edge damage occursin punched specimens from alloy with lower austenite stability.

TABLE 8 Tensile Properties of Punched vs EDM Cut Specimens from SelectedAlloys Yield Ultimate Tensile Cutting Strength Tensile StrengthElongation Alloy Method (MPa) (MPa) (%) Alloy 1 EDM Cut 392 1310 46.7397 1318 45.1 400 1304 49.7 Punched 431 699 9.3 430 680 8.1 422 656 6.9Alloy 2 EDM Cut 434 1213 46.4 452 1207 46.8 444 1199 49.1 Punched 491823 14.4 518 792 11.3 508 796 11.9 Alloy 9 EDM Cut 468 1166 56.1 4801177 52.4 475 1169 56.9 Punched 508 1018 29.2 507 1007 28.6 490 945 23.3Alloy 11 EDM Cut 474 1115 64.4 464 1165 62.5 495 1127 62.7 Punched 503924 24.6 508 964 28.0 490 921 25.7 Alloy 12 EDM Cut 481 1094 54.4 4791128 64.7 495 1126 62.4 Punched 521 954 27.1 468 978 30.7 506 975 31.2Alloy 13 EDM Cut 454 1444 39.5 450 1455 38.7 Punched 486 620 5.0 469 5996.3 483 616 4.5 Alloy 14 EDM Cut 484 1170 58.7 489 1182 61.2 468 118859.0 Punched 536 846 17.0 480 816 18.4 563 870 17.5 Alloy 18 EDM Cut 4451505 37.8 422 1494 37.5 Punched 478 579 2.4 469 561 2.6 463 582 2.9Alloy 21 EDM Cut 464 1210 57.6 499 1244 49.0 516 1220 54.5 Punched 527801 11.3 511 806 12.6 545 860 15.2 Alloy 24 EDM Cut 440 1166 31.0 4431167 32.0 455 1176 31.0 Punched 496 696 5.0 463 688 5.0 440 684 4.0Alloy 25 EDM Cut 474 1183 15.8 470 1204 17.0 485 1223 17.4 Punched 503589 2.1 517 579 0.8 497 583 2.1 Alloy 26 EDM Cut 735 1133 20.8 742 110919.0 Punched 722 898 3.4 747 894 2.9 764 894 3.1 Alloy 27 EDM Cut 5371329 19.3 513 1323 21.4 480 1341 20.8 Punched 563 624 4.3 568 614 3.3539 637 4.3 Alloy 34 EDM Cut 460 1209 54.7 441 1199 54.1 475 1216 52.9Punched 489 828 15.4 486 811 14.6 499 813 14.8 Alloy 35 EDM Cut 431 119650.6 437 1186 52.0 420 1172 54.7 Punched 471 826 19.9 452 828 19.7 482854 19.7

TABLE 9 Tensile Elongation in Specimens with Different Cutting MethodsAverage Tensile Loss In Tensile Elongation (%) Elongation (E2/E1) AlloyEDM Cut (E1) Punched (E2) Min Max Alloy 1 47.2 8.1 0.14 0.21 Alloy 247.4 12.5 0.23 0.31 Alloy 9 55.1 27.0 0.41 0.56 Alloy 11 63.2 26.1 0.380.45 Alloy 12 60.5 29.7 0.42 0.57 Alloy 13 39.1 5.2 0.11 0.16 Alloy 1459.7 17.7 0.28 0.31 Alloy 18 37.6 2.6 0.06 0.08 Alloy 21 53.7 13.0 0.200.31 Alloy 24 31.3 4.7 0.13 0.16 Alloy 25 16.7 1.7 0.05 0.13 Alloy 2631.3 4.7 0.14 0.18 Alloy 27 20.5 4.0 0.15 0.22 Alloy 34 53.9 14.9 0.270.29 Alloy 35 52.4 19.8 0.36 0.39

As can be seen from Table 30, EDM cutting is considered to berepresentative of the optimal mechanical properties of the identifiedalloys, without a sheared edge, and which were processed to the point ofassuming Structure #4 (Recrystallized Modal Structure). Accordingly,samples having a sheared edge due to punching indicate a significantdrop in ductility as reflected by tensile elongation measurements of thepunched samples having the ASTM E8 geometry. For Alloy 1, tensileelongation is initially 47.2% and then drops to 8.1%, a drop itself of82.8%%. The drop in ductility from the punched to the EDM cut (E2/E1)varies from 0.57 to 0.05.

The edge status after punching and EDM cutting was analyzed by SEM usingan EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMTInc. The typical appearance of the specimen edge after EDM cutting isshown for Alloy 1 in FIG. 36a . The EDM cutting method minimizes thedamage of a cut edge allowing the tensile properties of the material tobe measured without any deleterious edge effects. In wire-EDM cutting,material is removed from the edge by a series of rapidly recurringcurrent discharges/sparks and by this route an edge is formed withoutsubstantial deformation or edge damage. The appearance of the shearededge after punching is shown in FIG. 36b . A significant damage of theedge occurs in a fracture zone that undergoes severe deformation duringpunching leading to structural transformation in the shear affected zoneinto a Refined High Strength Nanomodal Structure (FIG. 37b ) withlimited ductility while Recrystallized Modal Structure was observed nearEDM cut edge (FIG. 37a ).

This Case Example demonstrates that in a case of wire-EDM cuttingtensile properties are measured at relative higher level as compared tothat after punching. In contrast to EDM cutting, punching of the tensilespecimens creates a significant edge damage which results in tensileproperty decrease. Relative excessive plastic deformation of the sheetalloys herein during punching leads to structural transformation to aRefined High Strength Nanomodal Structure (Structure #5, FIG. 1B) withreduced ductility leading to premature cracking at the edge andrelatively lower properties (e.g. reduction in elongation and tensilestrength). The magnitude of this drop in tensile properties has alsobeen observed to depend on the alloy chemistry in correlation withaustenite stability.

Case Example #7 Punched Edge vs EDM Cut Tensile Properties and Recovery

Slabs with thickness of 50 mm were laboratory cast from selected alloyslisted in Table 31 according to the atomic ratios provided in Table 2and laboratory processed by hot rolling, cold rolling and annealing at850° C. for 10 min as described herein. Resultant sheet from each alloywith final thickness of 1.2 mm and Recrystallized Modal Structure(Structure #4, FIG. 1B) was used to demonstrate edge damage recovery byannealing of punched tensile specimens. In the broad context of thepresent invention, annealing may be achieved by various methods,including but not limited to furnace heat treatment, induction heattreatment and/or laser heat treatment.

Tensile specimens in the ASTM E8 geometry were prepared using both wireEDM cutting and punching. Part of punched tensile specimens was then putthrough a recovery anneal of 850° C. for 10 minutes, followed by an aircool, to confirm the ability to recover properties lost by punching andshearing damage. Tensile properties were measured on an Instron 5984mechanical testing frame using Instron's Bluehill control software. Alltests were conducted at room temperature, with the bottom grip fixed andthe top grip set to travel upwards at a rate of 0.012 mm/s. Strain datawas collected using Instron's Advanced Video Extensometer. Tensiletesting results are provided in Table 31 and illustrated in FIG. 38 forselected alloys showing a substantial mechanical property recovery inpunched samples after annealing.

For example, in the case of Alloy 1 indicated, when EDM cut into atensile testing sample, a tensile elongation average value is about47.2%. As noted above, when punched and therefore containing a shearededge, the tensile testing of the sample with such edge indicated asignificant drop in such elongation values, i.e. an average value ofonly about 8.1% due to Mechanism #4 and formation of Refined HighStrength Nanomodal Structure (Structure #5, FIG. 1B), which whilepresent largely at the edge section where shearing occurred, isnonetheless reflected in the bulk property measurements in tensiletesting. However, upon annealing, which is representative of Mechanism#3 in FIG. 1B and conversion to Structure #4 (Recrystallized ModalStructure, FIG. 1B), the tensile elongation properties are restored. Inthe case of Alloy 1, the tensile elongation are brought back to anaverage value of about 46.2%. Example tensile stress-strain curves forpunched specimens from Alloy 1 with and without annealing are shown inFIG. 39. In Table 32, a summary of the average tensile properties andthe average lost and gained in tensile elongation is provided. Note thatthe individual losses and gains are a larger spread than the averagelosses. Accordingly, in the context of the present disclosure, thealloys herein, having an initial value of tensile elongation (E₁) whensheared, may indicate a drop in elongation properties to a value of E₂,wherein E₂=(0.0.57 to 0.05)(E₁). Then, upon application of Mechanism #3,which is preferably accomplished by heating/annealing at a temperaturerange of 450° C. up to the T_(m) depending on alloy chemistry, the valueof E₂ is recovered to an elongation value E₃=(0.48 to 1.21)(E₁).

TABLE 10 Tensile Properties of Punched and Annealed Specimens fromSelected Alloys Yield Ultimate Tensile Cutting Strength Tensile StrengthElongation Alloy Method (MPa) (MPa) (%) Alloy 1 EDM Cut 392 1310 46.7397 1318 45.1 400 1304 49.7 Punched 431 699 9.3 430 680 8.1 422 656 6.9Punched & 364 1305 43.6 Annealed 364 1315 47.6 370 1305 47.3 Alloy 2 EDMCut 434 1213 46.4 452 1207 46.8 444 1199 49.1 Punched 491 823 14.4 518792 11.3 508 796 11.9 Punched & 432 1205 50.4 Annealed 426 1191 50.7 4381188 49.3 Alloy 9 EDM Cut 468 1166 56.1 480 1177 52.4 475 1169 56.9Punched 508 1018 29.2 507 1007 28.6 490 945 23.3 Punched & 411 1166 59.0Annealed 409 1174 52.7 418 1181 55.6 Alloy 11 EDM Cut 474 1115 64.4 4641165 62.5 495 1127 62.7 Punched 503 924 24.6 508 964 28.0 490 921 25.7Punched & 425 1128 64.5 Annealed 429 1117 57.1 423 1140 54.3 Alloy 12EDM Cut 481 1094 54.4 479 1128 64.7 495 1126 62.4 Punched 521 954 27.1468 978 30.7 506 975 31.2 Punched & 419 1086 65.7 Annealed 423 1085 63.0415 1100 53.8 Alloy 13 EDM Cut 454 1444 39.5 450 1455 38.7 Punched 486620 5.0 469 599 6.3 483 616 4.5 Punched & 397 1432 41.4 Annealed 3971437 37.4 404 1439 40.3 Alloy 14 EDM Cut 484 1170 58.7 489 1182 61.2 4681188 59.0 Punched 536 846 17.0 480 816 18.4 563 870 17.5 Punched & 4231163 58.3 Annealed 412 1168 55.9 415 1177 51.5 Alloy 18 EDM Cut 445 150537.8 422 1494 37.5 Punched 478 579 2.4 469 561 2.6 463 582 2.9 Punched &398 1506 36.3 Annealed 400 1502 40.3 392 1518 35.4 Alloy 21 EDM Cut 4641210 57.6 499 1244 49.0 516 1220 54.5 Punched 527 801 11.3 511 806 12.6545 860 15.2 Punched & 409 1195 47.7 Annealed 418 1214 53.8 403 119451.8 Alloy 24 EDM Cut 440 1166 31.0 443 1167 32.0 455 1176 31.0 Punched496 696 5.0 463 688 5.0 440 684 4.0 Punched & 559 1100 22.3 Annealed 5811113 22.0 561 1100 22.3 Alloy 25 EDM Cut 474 1183 15.8 470 1204 17.0 4851223 17.4 Punched 503 589 2.1 517 579 0.8 497 583 2.1 Punched & 457 114315.4 Annealed 477 1159 14.6 423 1178 16.3 Alloy 26 EDM Cut 735 1133 20.8742 1109 19.0 Punched 722 898 3.4 747 894 2.9 764 894 3.1 Punched & 7151112 18.8 Annealed 713 1098 17.8 709 931 10.0 Alloy 27 EDM Cut 537 132919.3 513 1323 21.4 480 1341 20.8 Punched 563 624 4.3 568 614 3.3 539 6374.3 Punched & 505 1324 19.7 Annealed 514 1325 20.0 539 1325 19.4 Alloy29 EDM Cut 460 1209 54.7 441 1199 54.1 475 1216 52.9 Punched 489 82815.4 486 811 14.6 499 813 14.8 Punched & 410 1204 53.9 Annealed 410 122053.2 408 1214 52.3 Alloy 32 EDM Cut 431 1196 50.6 437 1186 52.0 420 117254.7 Punched 471 826 19.9 452 828 19.7 482 854 19.7 Punched & 406 116958.1 Annealed 403 1170 51.4 405 1176 57.7

TABLE 32 Summary of Tensile Properties; Loss (E2/E1) and Gain (E3/E1)Loss In Tensile Gain in Tensile Elongation (E2/E1) Elongation (E3/E1)Alloy Min Max Min Max Alloy 1 0.14 0.21 0.88 1.06 Alloy 2 0.23 0.31 1.001.09 Alloy 9 0.41 0.56 0.93 1.13 Alloy 11 0.38 0.45 0.84 1.03 Alloy 120.42 0.57 0.83 1.21 Alloy 13 0.11 0.16 0.95 1.07 Alloy 14 0.28 0.31 0.840.99 Alloy 18 0.06 0.08 0.94 1.07 Alloy 21 0.20 0.31 0.83 1.10 Alloy 240.13 0.16 0.69 0.72 Alloy 25 0.05 0.13 0.89 1.03 Alloy 26 0.14 0.18 0.480.99 Alloy 27 0.15 0.22 0.91 1.04 Alloy 29 0.27 0.29 0.97 1.02 Alloy 320.36 0.39 0.94 1.15

Punching of tensile specimens results in edge damage and lowering thetensile properties of the material. Plastic deformation of the sheetalloys herein during punching leads to structural transformation to aRefined High Strength Nanomodal Structure (Structure #5, FIG. 1B) withreduced ductility leading to premature cracking at the edge andrelatively lower properties (e.g. reduction in elongation and tensilestrength). This Case Example demonstrates that due to the uniquestructural reversibility, the edge damage in the alloys listed in Table2 is substantially recoverable by annealing leading back toRecrystallized Modal Structure (Structure #4, FIG. 1B) formation withfull or partial property restoration that depends on alloy chemistry andprocessing. For example, as exemplified by Alloy 1, punching andshearing and creating a sheared edge is observed to reduce tensilestrength from an average of about 1310 MPa (an EDM cut sample without asheared/damaged edge) to an average value of 678 MPa, a drop of between45 to 50%. Upon annealing, tensile strength recovers to an average valueof about 1308 MPa, which is in the range of greater than or equal to 95%of the original value of 1310 MPa. Similarly, tensile elongation isinitially at an average of about 47.1%, dropping to an average value of8.1%, a decrease of up to about 80 to 85%, and upon annealing andundergoing what is shown in FIG. 1B as Mechanism #3, tensile elongationrecovers to an average value of 46.1%, a recovery of greater than orequal to 90% of the value of the elongation value of 47.1%.

Case Example #8 Temperature Effect on Recovery and Recrystallization

Slabs with thickness of 50 mm were laboratory cast from Alloy 1 andlaboratory processed by hot rolling down to thickness of 2 mm and coldrolling with reduction of approximately 40%. Tensile specimens in theASTM E8 geometry were prepared by wire EDM cut from cold rolled sheet.Part of tensile specimens was annealed for 10 minutes at differenttemperatures in a range from 450 to 850° C., followed by an air cool.Tensile properties were measured on an Instron 5984 mechanical testingframe using Instron's Bluehill control software. All tests wereconducted at room temperature, with the bottom grip fixed and the topgrip set to travel upwards at a rate of 0.012 mm/s. Strain data wascollected using Instron's Advanced Video Extensometer. Tensile testingresults are shown in FIG. 40 demonstrating a transition in deformationbehavior depending on annealing temperature. During the process of coldrolling, the Dynamic Nanophase Strengthening (Mechanism #2, FIG. 1A) orthe Nanophase Refinement & Strengthening (Mechanism #4, FIG. 1B) occurswhich involves, once the yield strength is exceeded with increasingstrain, the continuous transformation of austenite to ferrite plus oneor more types of nanoscale hexagonal phases. Concurrent with thistransformation, deformation by dislocation mechanisms also occurs in thematrix grains prior to and after transformation. The result is thechange in the microstructure from the Nanomodal Structure (Structure #2,FIG. 1A) to the High Strength Nanomodal Structure (Structure #3, FIG.1A) or from the Recrystallized Modal Structure (Structure #4, FIG. 1B)to the Refined High Strength Nanomodal Structure (Structure #5, FIG.1B). The structure and property changes occurring during colddeformation can be reversed at various degrees by annealing depending onannealing parameters as seen in the tensile curves of FIG. 40A. In FIG.40B, the corresponding yield strength from the tensile curves areprovided as a function of the heat treatment temperature. The yieldstrength after cold rolling with no anneal is measured at 1141 MPa. Asshown, depending on how the material is annealed which may includepartial and full recovery and partial and full recrystallization theyield strength can be varied widely from 1372 MPa at the 500° C. annealdown to 458 MPa at the 850° C. anneal.

To show the microstructural recovery in accordance to the tensileproperty upon annealing, TEM studies were conducted on selected samplesthat were annealed at different temperatures. For comparison, coldrolled sheet was included as a baseline herein. Laboratory cast Alloy 1slab of 50 mm thick was used, and the slab was hot rolled at 1250° C. bytwo-step of 80.8% and 78.3% to approx. 2 mm thick, then cold rolled by37% to sheet of 1.2 mm thick. The cold rolled sheet was annealed at 450°C., 600° C., 650° C. and 700° C. respectively for 10 minutes. FIG. 41shows the microstructure of as-cold rolled Alloy 1 sample. It can beseen that typical High Strength Nanomodal Structure is formed after coldrolling, in which high density of dislocations are generated along withthe presence of strong texture. Annealing at 450° C. for 10 min does notlead to recrystallization and formation of the High Strength NanomodalStructure, as the microstructure remains similar to that of the coldrolled structure and the rolling texture remains unchanged (FIG. 42).When the cold rolled sample is annealed at 600° C. for 10 min, TEManalysis shows very small isolated grains, a sign of the beginning ofrecrystallization. As shown in FIG. 43, isolated grains of 100 nm or soare produced after the annealing, while areas of deformed structure withdislocation networks are also present. Annealing at 650° C. for 10 minshows larger recrystallized grains suggesting the progress ofrecrystallization. Although the fraction of deformed area is reduced,the deformed structure continues to be seen, as shown in FIG. 44.Annealing at 700° C. 10 min shows larger and cleaner recrystallizedgrains, as displayed by FIG. 45. Selected electron diffraction showsthat these recrystallized grains are of the austenite phase. The area ofdeformed structure is smaller compared to the samples annealed at lowertemperature. Survey over the entire sample suggests that approx. 10% to20% area is occupied by the deformed structure. The progress ofrecrystallization revealed by TEM in the samples annealed at lowertemperature to higher temperature corresponds excellently to the changeof tensile properties shown in FIG. 40. These low temperature annealedsamples (such as below 600° C.) maintain predominantly the High StrengthNanomodal Structure, leading to the reduced ductility. Therecrystallized sample (such as at 700° C.) recovers majority of theelongation, compared to the fully recrystallized sample at 850° C. Theannealing in between these temperatures partially recovers theductility.

One reason behind the difference in recovery and transition indeformation behavior is illustrated by the model TTT diagram in FIG. 46.As described previously, the very fine/nanoscale grains of ferriteformed during cold working recrystallize into austenite during annealingand some fraction of the nanoprecipitates re-dissolve. Concurrently, theeffect of the strain hardening is eliminated with dislocation networksand tangles, twin boundaries, and small angle boundaries beingannihilated by various known mechanisms. As shown by the heating curve Aof the model temperature, time transformation (TTT) diagram in FIG. 46,at low temperatures (particularly below 650° C. for Alloy 1), onlyrecovery may occur without recrystallization (i.e. recovery being areference to a reduction in dislocation density).

In other words, in the broad context of the present invention, theeffect of shearing and formation of a sheared edge, and its associatednegative influence on mechanical properties, can be at least partiallyrecovered at temperatures of 450° C. up to 650° C. as shown in FIG. 46.In addition, at 650° C. and up to below Tm of the alloy,recrystallization can occur, which also contributes to restoringmechanical strength lost due to the formation of a sheared edge.

Accordingly, this Case Example demonstrates that upon deformation duringcold rolling, concurrent processes occur involving dynamic strainhardening and phase transformation through unique Mechanisms #2 or #3(FIG. 1A) along with dislocation based mechanisms. Upon heating, themicrostructure can be reversed into a Recrystallized Modal Structure(Structure #4, FIG. 1B). However, at low temperatures, this reversingprocess may not occur when only dislocation recovery takes place. Thus,due to the unique mechanisms of the alloys in Table 2, various externalheat treatments can be used to heal the edge damage frompunching/stamping.

Case Example #9 Temperature Effect of Punched Edge Recovery

Slabs with thickness of 50 mm were laboratory cast from selected alloyslisted in Table 33 according to the atomic ratios provided in Table 2and laboratory processed by hot rolling, cold rolling and annealing at850° C. for 10 min as described in Main Body section of currentapplication. Resultant sheet from each alloy with final thickness of 1.2mm and Recrystallized Modal Structure (Structure #4, FIG. 1B) was usedto demonstrate punched edge damage recovery after annealing as afunction of temperature.

Tensile specimens in the ASTM E8 geometry were prepared by punching. Apart of punched tensile specimens from selected alloys was then putthrough a recovery anneal for 10 minutes at different temperatures in arange from 450 to 850° C., followed by an air cool. Tensile propertieswere measured on an Instron 5984 mechanical testing frame usingInstron's Bluehill control software. All tests were conducted at roomtemperature, with the bottom grip fixed and the top grip set to travelupwards at a rate of 0.012 mm/s. Strain data was collected usingInstron's Advanced Video Extensometer.

Tensile testing results are shown in Table 32 and in FIG. 47. As it canbe seen, full or nearly full property recovery achieved after annealingat temperatures at 650° C. and higher, suggesting that the structure isfully or near fully recrystallized (i.e. change in structure fromStructure #5 to Structure #4 in FIG. 1B) in the damaged edges afterpunching. For example, the level of recrystallization at the damagededge is contemplated to be at a level of greater than or equal to 90%when annealing temperatures are in the range of 650° C. up to T_(m).Lower annealing temperature (e.g. temperatures below 650° C. does notresult in full recrystallization and leads to partial recovery (i.e.decrease in dislocation density) as described in Case Example #8 andillustrated in FIG. 6.

Microstructural changes in Alloy 1 at the shear edge as a result of thepunching and annealing at different temperatures were examined by SEM.Cross section samples were cut from ASTM E8 punched tensile specimensnear the sheared edge in as-punched condition and after annealing at650° C. and 700° C. as shown in FIG. 48.

For SEM study, the cross section samples were ground on SiC abrasivepapers with reduced grit size, and then polished progressively withdiamond media paste down to 1 μm. The final polishing was done with 0.02μm grit SiO2 solution. Microstructures were examined by SEM using anEVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMTInc.

FIG. 49 shows the backscattered SEM images of the microstructure at theedge in the as-punched condition. It can be seen that the microstructureis deformed and transformed in the shear affected zone (i.e., thetriangle with white contrast close to the edge) in contrast to therecrystallized microstructure in the area away from the shear affectedzone. Similar to tensile deformation, the deformation in the shearaffected zone caused by punching creates Refined High Strength NanomodalStructure (Structure #5, FIG. 1B) through Nanophase Refinement &Strengthening mechanism. However, annealing recovers the tensileproperties of punched ASTM E8 specimens, which are related to themicrostructure change in the shear affected zone during annealing. FIG.50 shows the microstructure of the sample annealed at 650° C. for 10minutes. Compared to the as-punched sample, the shear affected zonebecomes smaller with less contrast suggesting that the microstructure inthe shear affected zone evolves toward that in the center of the sample.A high magnification SEM image shows that some very small grains arenucleated, but recrystallization does not take place massively acrossthe shear affected zone. It is likely that the recrystallization is inthe early stage with most of the dislocations annihilated. Although thestructure is not fully recrystallized, the tensile property issubstantially recovered (Table 32 and FIG. 47a ). Annealing at 700° C.for 10 minutes leads to full recrystallization of the shear affectedzone. As shown in FIG. 51, the contrast in shear affected zonesignificantly decreased. High magnification image shows that equiaxedgrains with clear grain boundaries are formed in the shear affectedzone, indicating full recrystallization. The grain size is smaller thanthat in the center of sample. Note that the grains in the center areresulted from recrystallization after annealing at 850° C. for 10minutes before punching of specimens. With the shear affected zone fullyrecrystallized, the tensile properties are fully recovered, as shown inTable 32 and FIG. 47 a.

Punching of tensile specimens result in edge damage lowering the tensileproperties of the material. Plastic deformation of the sheet alloysherein during punching leads to structural transformation to a RefinedHigh Strength Nanomodal Structure (Structure #5, FIG. 1B) with reducedductility leading to premature cracking at the edge. This Case Exampledemonstrates that this edge damage is partially/fully recoverable bydifferent anneals over a wide range of industrial temperatures.

TABLE 33 Tensile Properties after Punching and Annealing at DifferentTemperatures Anneal Yield Ultimate Tensile Temperature Strength TensileStrength Elongation Alloy (° C.) (MPa) (MPa) (%) Alloy 1 As Punched 494798 12.6 487 829 14.3 474 792 15.3 450 481 937 21.5 469 934 20.9 485 85219.3 600 464 1055 27.3 472 1103 30.5 453 984 23.7 650 442 1281 51.5 4541270 45.4 445 1264 51.1 700 436 1255 50.1 442 1277 52.1 462 1298 51.6850 407 1248 52.0 406 1260 47.8 412 1258 48.5 Alloy 9 As Punched 5081018 29.2 507 1007 28.6 490 945 23.3 600 461 992 28.5 462 942 24.8 471968 25.6 650 460 1055 33.0 470 1166 48.3 473 1177 49.3 700 457 1208 57.5455 1169 50.3 454 1171 61.6 850 411 1166 59.0 409 1174 52.7 418 118155.6 Alloy 12 As Punched 521 954 27.1 468 978 30.7 506 975 31.2 600 4621067 44.9 446 1013 41.3 471 1053 41.1 650 452 1093 61.5 449 1126 57.8505 1123 55.4 700 480 1112 59.6 460 1117 61.8 468 1096 61.5 850 419 108665.7 423 1085 63.0 415 1100 53.8

Case Example #10 Effect of Punching Speed on Punched Edge PropertyReversibility

Slabs with thickness of 50 mm were laboratory cast from selected alloyslisted in Table 34 according to the atomic ratios provided in Table 2and laboratory processed by hot rolling, cold rolling and annealing at850° C. for 10 min as described herein. Resultant sheet from each alloywith final thickness of 1.2 mm and Recrystallized Modal Structure(Structure #4, FIG. 1B) was used to demonstrate edge damage recovery asa function of punching speed.

Tensile specimens in the ASTM E8 geometry were prepared by punching atthree different speeds of 28 mm/s, 114 mm/s, and 228 mm/s. Wire EDM cutspecimens from the same materials were used for the reference. A part ofpunched tensile specimens from selected alloys was then put through arecovery anneal for 10 minutes at 850° C., followed by an air cool.Tensile properties were measured on an Instron 5984 mechanical testingframe using Instron's Bluehill control software. All tests wereconducted at room temperature, with the bottom grip fixed and the topgrip set to travel upwards at a rate of 0.012 mm/s. Strain data wascollected using Instron's Advanced Video Extensometer. Tensile testingresults are listed in Table 34 and tensile properties as a function ofpunching speed for selected alloys are illustrated in FIG. 52. It isseen that tensile properties drop significantly in the punched samplesas compared to that for wire EDM cut. Punching speed increase from 28mm/s to 228 mm/s leads to increase in properties of all three selectedalloys. The localized heat generation during punching a hole or shearingan edge is known to increase with increasing punching velocity and mightbe a factor in edge damage recovery in specimens punched at higherspeed. Note that heat alone will not cause edge damage recovery but willbe enabled by the materials response to the heat generated. Thisdifference in response for the alloys contained in Table 2 in thisapplication to commercial steel samples is clearly illustrated in CaseExamples 15 and 17.

TABLE 34 Tensile Properties of Specimens Punched at Different Speed vsEDM Cut Sample Yield Tensile Tensile Preparation Strength StrengthElongation Alloy Method (MPa) (MPa) (%) Alloy 1 EDM 459 1255 51.2 4431271 46.4 441 1248 52.7 453 1251 55.0 467 1259 51.3 228 mm/s 474 95221.8 Punched 498 941 21.6 493 956 21.6 114 mm/s 494 798 13.4 Punched 487829 15.1 474 792 14.1 28 mm/s 464 770 12.8 Punched 479 797 13.7 465 75512.1 Alloy 9 EDM 468 1166 56.1 480 1177 52.4 475 1169 56.9 228 mm/s 5001067 35.1 Punched 493 999 28.8 470 1042 31.8 114 mm/s 508 1018 29.2Punched 507 1007 28.6 490 945 23.3 28 mm/s 473 851 19.7 Punched 472 84116.4 494 846 18.9 Alloy 12 EDM 481 1094 54.4 479 1128 64.7 495 1126 62.4228 mm/s 495 1124 53.8 Punched 484 1123 53.0 114 mm/s 521 954 27.1Punched 468 978 30.7 506 975 31.2 28 mm/s 488 912 23.6 Punched 472 90021.7 507 928 22.9

This Case Example demonstrates that punching speed can have asignificant effect on the resulting tensile properties in steel alloysherein. Localized heat generation at punching might be a factor inrecovery of the structure near the edge leading to property improvement.

Case Example #11 Edge Structure Transformation During Hole Punching andExpansion

Slabs with thickness of 50 mm were laboratory cast from Alloy 1 andlaboratory processed by hot rolling, cold rolling and annealing at 850°C. for 10 min as described herein. Resultant sheet with final thicknessof 1.2 mm and Recrystallized Modal Structure (Structure #4, FIG. 1B) wasused for hole expansion ratio (HER) tests.

Specimens for testing with a size of 89×89 mm were wire EDM cut from thesheet. The hole with 10 mm diameter was cut in the middle of specimensby utilizing two methods: punching and drilling with edge milling. Thehole punching was done on an Instron Model 5985 Universal Testing Systemusing a fixed speed of 0.25 mm/s with 16% clearance. Hole expansionratio (HER) testing was performed on the SP-225 hydraulic press andconsisted of slowly raising the conical punch that uniformly expandedthe hole radially outward. A digital image camera system was focused onthe conical punch and the edge of the hole was monitored for evidence ofcrack formation and propagation. The initial diameter of the hole wasmeasured twice with calipers, measurements were taken at 90° incrementsand averaged to get the initial hole diameter. The conical punch wasraised continuously until a crack was observed propagating through thespecimen thickness. At that point the test was stopped and the holeexpansion ratio was calculated as a percentage of the initial holediameter measured before the start of the test. After expansion fourdiameter measurements were taken using calipers every 45° and averagedto account for any asymmetry of the hole due to cracking.

Results of HER testing are shown in FIG. 53 demonstrating asignificantly lower value for the sample when the hole was prepared bypunching as compared to milling: 5.1% HER vs 73.6% HER, respectively.Samples were cut from both tested samples as shown in FIG. 54 for SEManalysis and microhardness measurements.

Microhardness was measured for Alloy 1 at all relevant stages of thehole expansion process. Microhardness measurements were taken alongcross sections of sheet samples in the annealed (before punching and HERtesting), as-punched, and HER tested conditions. Microhardness was alsomeasured in cold rolled sheet from Alloy 1 for reference. Measurementprofiles started at an 80 micron distance from the edge of the sample,with an additional measurement taken every 120 microns until 10 suchmeasurements were taken. After that point, further measurements weretaken every 500 microns, until at least 5 mm of total sample length hadbeen measured. A schematic illustration of microhardness measurementlocations in HER tested samples is shown in FIG. 55. SEM images of thepunched and HER tested samples after microhardness measurements areshown in FIG. 56.

As shown in FIG. 57, the punching process creates a transformed zone ofapproximately 500 microns immediately adjacent to the punched edge, withthe material closest to the punched edge either fully or near-fullytransformed, as evidenced by the hardness approaching that observed inthe fully-transformed, 40% cold rolled material immediately next to thepunched edge. Microhardness profiles for each sample is presented inFIG. 58. As it can be seen, microhardness gradually increases towards ahole edge in the case of milled while in the case of punched holemicrohardness increase was observed in a very narrow area close to thehole edge. TEM samples were cut at the same distance in both cases asindicated in FIG. 58.

To prepare the TEM specimens, the HER test samples were first sectionedby wire EDM, and a piece with a portion of hole edge was thinned bygrinding with pads of reduced grit size. Further thinning to ˜60 μmthickness is done by polishing with 9 μm, 3 μm, and 1 μm diamondsuspension solution respectively. Discs of 3 mm in diameter were punchedfrom the foils near the edge of the hole and the final polishing wascompleted by electropolishing using a twin-jet polisher. The chemicalsolution used was a 30% Nitric acid mixed in Methanol base. In case ofinsufficient thin area for TEM observation, the TEM specimens may beion-milled using a Gatan Precision Ion Polishing System (PIPS). Theion-milling usually is done at 4.5 keV, and the inclination angle isreduced from 4° to 2° to open up the thin area. The TEM studies weredone using a JEOL 2100 high-resolution microscope operated at 200 kV.Since the location for TEM study is at the center of the disc, theobserved microstructure is approximately ˜1.5 mm from the edge of hole.

The initial microstructure of the Alloy 1 sheet before testing is shownon FIG. 59 representing Recrystallized Modal Structure (Structure #4,FIG. 1B). FIG. 60a shows the TEM micrograph of the microstructure in theHER test sample with punched hole after testing (HER=5.1%) in differentareas at the location of 1.5 mm from hole edge. It was found that mainlythe recrystallized microstructure remains in the sample (FIG. 60a ) withsmall amount of area with partially transformed “pockets” (FIG. 60b )indicating that limited volume (˜1500 μm deep) of the sample wasinvolved in deformation at HER testing. In the HER sample with milledhole (HER=73.6%), as shown in FIG. 61, there is a great amount ofdeformation in the sample as indicated by a large amount of transformed“pockets” and high density of dislocations (10⁸ to 10¹⁰ mm⁻²).

To analyze in more detail the reason causing the poor HER performance insamples with punched holes, Focused Ion Beam (FIB) technique wasutilized to make TEM specimens at the very edge of the punched hole. Asshown in FIG. 62, TEM specimen is cut at ˜10 μm from the edge. Toprepare TEM specimens by FIB, a thin layer of platinum is deposited onthe area to protect the specimen to be cut. A wedge specimen is then cutout and lifted by a tungsten needle. Further ion milling is performed tothin the specimen. Finally the thinned specimen is transferred andwelded to copper grid for TEM observation. FIG. 63 shows themicrostructure of the Alloy 1 sheet at the distance of ˜10 micron fromthe punched hole edge which is significantly refined and transformed ascompared to the microstructure in the Alloy 1 sheet before punching. Itsuggests that punching caused severe deformation at the hole edge suchthat Nanophase Refinement & Strengthening (Mechanism #4, FIG. 1B)occurred leading to formation of Refined High Strength NanomodalStructure (Structure #5, FIG. 1B) in the area close to the punched holeedge. This structure has relative lower ductility as compared toRecrystallized Modal Structure Table 1 resulting in premature crackingat the edge and low HER values. This Case Example demonstrates that thealloys in Table 2 exhibit the unique ability to transform from aRecrystallized Modal Structure (Structure #4, FIG. 1B) to a Refined HighStrength Nanomodal Structure (Structure #5, FIG. 1B) through theidentified Nanophase Refinement & Strengthening (Mechanism #4, FIG. 1B).The structural transformation occurring due to deformation at the holeedge at punching appears to be similar in nature to transformationoccurring during cold rolling deformation and that observed duringtensile testing deformation.

Case Example #12 HER Testing Results with and without Annealing

Slabs with thickness of 50 mm were laboratory cast from selected alloyslisted in Table 35 according to the atomic ratios provided in Table 2and laboratory processed by hot rolling, cold rolling and annealing at850° C. for 10 min as described herein. Resultant sheet with finalthickness of 1.2 mm and Recrystallized Modal Structure (Structure #4,FIG. 1B) was used for hole expansion ratio (HER) tests.

Test specimens of 89×89 mm were wire EDM cut from the sheet from largersections. A 10 mm diameter hole was made in the center of specimens bypunching on an Instron Model 5985 Universal Testing System using a fixedspeed of 0.25 mm/s at 16% punch to die clearance. Half of the preparedspecimens with punched holes were individually wrapped in stainlesssteel foil and annealed at 850° C. for 10 minutes before HER testing.Hole expansion ratio (HER) testing was performed on the SP-225 hydraulicpress and consisted of slowly raising the conical punch that uniformlyexpanded the hole radially outward. A digital image camera system wasfocused on the conical punch and the edge of the hole was monitored forevidence of crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

The results of the hole expansion ratio measurements on the specimenswith and without annealing after hole punching are shown in Table 35. Asshown in FIG. 64, FIG. 65, FIG. 66, FIG. 67 and FIG. 68 for Alloy 1,Alloy 9, Alloy 12, Alloy 13, and Alloy 17, respectively, the holeexpansion ratio measured with punched holes with annealing is generallygreater than in punched holes without annealing. The increase in holeexpansion ratio with annealing for the identified alloys hereintherefore leads to an increase in the actual HER of about 25% to 90%.

TABLE 35 Hole Expansion Ratio Results for Select Alloys With and WithoutAnnealing Punch Measured Hole Average Hole Clearance Expansion RatioExpansion Ratio Material Condition (%) (%) (%) Alloy 1 Without 16 3.003.20 Annealing 3.90 2.70 With 16 105.89 93.10 Annealing 81.32 92.11Alloy 9 Without 16 3.09 3.19 Annealing 3.19 3.29 With 16 78.52 87.84Annealing 97.60 87.40 Alloy 12 Without 16 4.61 4.91 Annealing 5.21 With16 69.11 77.60 Annealing 83.60 80.08 Alloy 13 Without 16 1.70 1.53Annealing 1.40 1.50 With 16 32.37 31.12 Annealing 29.00 32.00 Alloy 17Without 16 12.89 21.46 Annealing 28.70 22.80 With 16 104.21 103.74Annealing 80.42 126.58

This Case Example demonstrates that edge formability demonstrated duringHER testing can yield poor results due to edge damage during thepunching operation as a result of the unique mechanisms in the alloyslisted in Table 2. The fully post processed alloys exhibit very hightensile ductility as shown in Table 6 through Table 10 coupled with veryhigh strain hardening and resistance to necking until near failure.Thus, the material resists catastrophic failure to a great extent butduring punching, artificial catastrophic failure is forced to occur nearthe punched edge. Due to the unique reversibility of the identifiedmechanisms, this deleterious edge damage as a result of NanophaseRefinement & Strengthening (Mechanism #3, FIG. 1A) and structuraltransformation can be reversed by annealing resulting in high HERresults. Thus, high hole expansion ratio values can be obtained in acase of punching hole with following annealing and retaining exceptionalcombinations of tensile properties and the associated bulk formability.

In addition, it can be appreciated that the alloys herein that haveundergone the processing pathways to provide such alloys in the form ofStructure #4 (Recrystallized Modal Structure) will indicate, for a holethat is formed by shearing, and including a sheared edge, a first holeexpansion ratio (HER₁) and upon heating the alloy will have a secondhole expansion ratio (HER₂), wherein HER₂>HER₁.

More specifically, it can also be appreciated that the alloys hereinthat have undergone the processing pathways to provide such alloys withStructure #4 (Recrystallized Modal Structure) will indicate, for a holethat was placed in the alloy through methods (i.e. waterjet cutting,laser cutting, wire-edm, milling etc.) where the hole that is formedthat does not rely primarily on shearing, compared to punching a hole, afirst hole expansion ratio (HER₁) where such value may itself fall inthe range of 30 to 130%. However, when the same alloy includes a holeformed by shearing, a second hole expansion ratio is observed (HER₂)wherein HER₂=(0.01 to 0.30)(HER₁). However, if the alloy is then subjectto heat treatment herein, it is observed that HER₂ is recovered to aHER₃=(0.60 to 1.0) HER₁.

Case Example #13 Edge Condition Effect on Alloy Properties

Slabs with thickness of 50 mm were laboratory cast from Alloy 1according to the atomic ratios provided in Table 2 and laboratoryprocessed by hot rolling, cold rolling and annealing at 850° C. for 10min as described herein. Resultant sheet from Alloy 1 with finalthickness of 1.2 mm and Recrystallized Modal Structure (Structure #4,FIG. 1B) was used to demonstrate the effect that edge condition has onAlloy 1 tensile and hole expansion properties.

Tensile specimens of ASTM E8 geometry were created using two methods:Punching and wire EDM cutting. Punched tensile specimens were createdusing a commercial press. A subset of punched tensile specimens was heattreated at 850° C. for 10 minutes to create samples with a punched thenannealed edge condition.

Tensile properties of ASTM E8 specimens were measured on an Instron 5984mechanical testing frame using Instron's Bluehill 3 control software.All tests were conducted at room temperature, with the bottom grip fixedand the top grip set to travel upwards at a rate of 0.025 mm/s for thefirst 0.5% elongation, and at a rate of 0.125 mm/s after that point.Strain data was collected using Instron's Advanced Video Extensometer.Tensile properties of Alloy 1 with punched, EDM cut, and punched thenannealed edge conditions are shown in Table 36. Tensile properties ofAlloy 1 with different edge conditions are shown in FIG. 69.

TABLE 36 Tensile Properties of Alloy 1 with Different Edge ConditionsUltimate Tensile Tensile Edge Elongation Strength Condition (%) (MPa)Punched 12.6 798 14.3 829 15.3 792 EDM Cut 50.5 1252 51.2 1255 52.7 124855.0 1251 51.3 1259 50.5 1265 Punched 52.0 1248 Then 47.8 1260 Annealed48.5 1258

Specimens for hole expansion ratio testing with a size of 89×89 mm werewire EDM cut from the sheet. The holes with 10 mm diameter were preparedby two methods: punching and cutting by wire EDM. The punched holes with10 mm diameter were created by punching at 0.25 mm/s on an Instron 5985Universal Testing System with a 16% punch clearance and with using theflat punch profile geometry. A subset of punched samples for holeexpansion testing were annealed with an 850° C. for 10 minutes heattreatment after punching.

Hole expansion ratio (HER) testing was performed on the SP-225 hydraulicpress and consisted of slowly raising the conical punch that uniformlyexpanded the hole radially outward. A digital image camera system wasfocused on the conical punch and the edge of the hole was monitored forevidence of crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

Hole expansion ratio testing results are shown in Table 37. An averagehole expansion ratio value for each edge condition is also shown. Theaverage hole expansion ratio for each edge condition is plotted in FIG.70. It can be seen that for samples with EDM cut and punched thenannealed edge conditions the edge formability (i.e. HER response) isexcellent, whereas samples with holes in the punched edge condition haveconsiderably lower edge formability.

TABLE 37 Hole Expansion Ratio of Alloy 1 with Different Edge ConditionsMeasured Average Hole Hole Expansion Expansion Edge Ratio RatioCondition (%) (%) Punched 3.00 3.20 3.90 2.70 EDM Cut 92.88 82.43 67.9486.47 Punched 105.90 93.10 Then 81.30 Annealed 92.10

This Case Example demonstrates that the edge condition of Alloy 1 has adistinct effect on the tensile properties and edge formability (i.e. HERresponse). Tensile samples tested with punched edge condition havediminished properties when compared to both wire EDM cut and punchedafter subsequent annealing. Samples having the punched edge conditionhave hole expansion ratios averaging 3.20%, whereas EDM cut and punchedthen annealed edge conditions have hole expansion ratios of 82.43% and93.10%, respectively. Comparison of edge conditions also demonstratesthat damage associated with edge creation (i.e. via punching) has anon-trivial effect on the edge formability of the alloys herein.

Case Example #14 HER Results as a Function of Hole Punching Speed

Slabs with thickness of 50 mm were laboratory cast from selected alloyslisted in Table 38 according to the atomic ratios provided in Table 2and laboratory processed by hot rolling, cold rolling and annealing at850° C. for 10 min as described herein. Resultant sheet from each alloywith final thickness of 1.2 mm and Recrystallized Modal Structure(Structure #4, FIG. 1B) were used to demonstrate an effect of holepunching speed on HER results.

Specimens for testing with a size of 89×89 mm were wire EDM cut from thesheet. The holes with 10 mm diameter were punched at different speeds ontwo different machines but all of the specimens were punched with a 16%punch clearance and with the same punch profile geometry. The low speedpunched holes (0.25 mm/s, 8 mm/s) were punched using an Instron 5985Universal Testing System and the high speed punched holes (28 mm/s, 114mm/s, 228 mm/s) were punched on a commercial punch press. All holes werepunched using a flat punch geometry.

Hole expansion ratio (HER) testing was performed on the SP-225 hydraulicpress and consisted of slowly raising the conical punch that uniformlyexpanded the hole radially outward. A digital image camera system wasfocused on the conical punch and the edge of the hole was monitored forevidence of crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

Hole expansion ratio values for tests are shown in Table 37. An averagehole expansion value is shown for each speed and alloy tested at 16%punch clearance. The average hole expansion ratio as a function of punchspeed is shown in FIG. 71, FIG. 72 and FIG. 73 for Alloy 1, Alloy 9, andAlloy 12, respectively. It can be seen that as punch speed increases,all alloys tested had a positive edge formability response, asdemonstrated by an increase in hole expansion ratio. The reason for thisincrease is believed to be related to the following effects. With higherpunch speed, the amount of heat generated at the sheared edge isexpected to increase and the localized temperature spike may result inan annealing effect (i.e. in-situ annealing). Alternatively, withincreasing punch speed, there may be a reduced amount of materialtransforming from the Recrystallized Modal Structure (i.e. Structure #4in FIG. 1B) to the Refined High Strength Nanomodal Structure (i.e.Structure #5 in FIG. 1B). Concurrently, the amount of Refined HighStrength Nanomodal Structure (i.e. Structure #5 in FIG. 1B) may bereduced due to the temperature spike enabling localizedrecrystallization (i.e. Mechanism #3 in FIG. 1B).

TABLE 38 Hole Expansion Ratio at Different Punch Speeds Measured AverageHole Hole Punch Expansion Expansion Speed Ratio Ratio Material (mm/s)(%) (%) Alloy 1 0.25 3.00 3.20 0.25 3.90 0.25 2.70 8 4.49 3.82 8 3.49 83.49 28 8.18 7.74 28 8.08 28 6.97 114 17.03 17.53 114 19.62 114 15.94228 20.44 21.70 228 21.24 228 23.41 Alloy 9 0.25 3.09 3.19 0.25 3.190.25 3.29 8 6.80 6.93 8 7.39 8 6.59 28 21.04 19.11 28 17.35 28 18.94 11424.80 24.29 114 19.74 114 28.34 228 26.00 30.57 228 35.16 228 30.55Alloy 12 0.25 4.61 4.91 0.25 5.21 8 7.62 11.28 8 14.61 8 11.62 28 29.3831.59 28 33.70 28 31.70 114 40.08 45.50 114 48.11 114 48.31 228 50.0049.36 228 40.56 228 57.51

This Case Example demonstrates a dependence of edge formability onpunching speed as measured by the hole expansion ratio. As punch speedincreases, the hole expansion ratio generally increases for the alloystested. With increased punching speed, the nature of the edge is changedsuch that improved edge formability (i.e. HER response) is achieved. Atpunching speeds greater than those measured, edge formability isexpected to continue improving towards even higher hole expansion ratiovalues.

Case Example #15 HER in DP980 as a Function of Hole Punching Speed

Commercially produced and processed Dual Phase 980 steel was purchasedand hole expansion ratio testing was performed. All specimens weretested in the as received (commercially processed) condition.

Specimens for testing with a size of 89×89 mm were wire EDM cut from thesheet. The holes with 10 mm diameter were punched at different speeds ontwo different machines but all of the specimens were punched with a 16%punch clearance and with the same punch profile geometry using acommercial punch press. The low speed punched holes (0.25 mm/s) werepunched using an Instron 5985 Universal Testing System and the highspeed punched holes (28 mm/s, 114 mm/s, 228 mm/s) were punched on acommercial punch press. All holes were punched using a flat punchgeometry.

Hole expansion ratio (HER) testing was performed on the SP-225 hydraulicpress and consisted of slowly raising the conical punch that uniformlyexpanded the hole radially outward. A digital image camera system wasfocused on the conical punch and the edge of the hole was monitored forevidence of crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

Values for hole expansion tests are shown in Table 39. The average holeexpansion value for each punching speed is also shown for commercialDual Phase 980 material at 16% punch clearance. The average holeexpansion value is plotted as a function of punching speed forcommercial Dual Phase 980 steel in FIG. 74.

TABLE 39 Hole Expansion Ratio of Dual Phase 980 Steel at Different PunchSpeeds Measured Average Hole Hole Punch Expansion Expansion Speed RatioRatio Material (mm/s) (%) (%) Commercial 0.25 23.55 22.45 Dual 0.2520.96 Phase 980 0.25 22.85 28 18.95 18.26 28 17.63 28 18.21 114 17.4020.09 114 23.66 114 19.22 228 27.21 23.83 228 24.30 228 19.98

This Case Example demonstrates that no edge performance effect based onpunch speed is measurable in Dual Phase 980 steel. For all punch speedsmeasured on Dual Phase 980 steel the edge performance (i.e. HERresponse) is consistently within the 21%±3% range, indicating that edgeperformance in conventional AHSS is not improved by punch speed asexpected since the unique structures and mechanisms present in thisapplication as for example in FIGS. 1a and 1b are not present.

Case Example #16: HER Results as a Function of Punch Design

Slabs with thickness of 50 mm were laboratory cast from Alloys 1, 9, and12 according to the atomic ratios provided in Table 2 and laboratoryprocessed by hot rolling, cold rolling and annealing at 850° C. for 10min as described herein. Resultant sheet from each alloy with finalthickness of 1.2 mm and Recrystallized Modal Structure (Structure #4,FIG. 1B) was used to demonstrate an effect of hole punching speed on HERresults.

Tested specimens of 89×89 mm were wire EDM cut from larger sections. A10 mm diameter hole was punched in the center of the specimen at threedifferent speeds, 28 mm/s, 114 mm/s, and 228 mm/s at 16% punch clearanceand with four punch profile geometries using a commercial punch press.These punch geometries used were flat, 6° tapered, 7° conical, andconical flat. Schematic drawings of the 6° tapered, 7° conical, andconical flat punch geometries are shown in FIG. 75.

Hole expansion ratio (HER) testing was performed on the SP-225 hydraulicpress and consisted of slowly raising the conical punch that uniformlyexpanded the hole radially outward. A digital image camera system wasfocused on the conical punch and the edge of the hole was monitored forevidence of crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

Hole expansion ratio data is included respectively in Table 40, Table41, and Table 42 for Alloy 1, Alloy 9, and Alloy 12 at four punchgeometries and at two different punch speeds. The average hole expansionvalues for Alloy 1, Alloy 9, and Alloy 12 are shown in FIG. 76, FIG. 77and FIG. 78, respectively. For all alloys tested, the 7° conical punchgeometry resulted in the largest or tied for the largest hole expansionratio compared to all other punch geometries. Increased punch speed isalso shown to improve the edge formability (i.e. HER response) for allpunch geometries. At increased punching speed with different punchgeometries, the alloys herein may be able to undergo some amount ofRecrystallization (Mechanism #3) as it is contemplated that there couldbe localized heating at the edge at such higher relative punch speeds,triggering Mechanism #3 and formation of some amount of Structure #4.

TABLE 40 Hole Expansion Ratio of Alloy 1 with Different Punch GeometriesMeasured Average Hole Hole Punch Expansion Expansion Punch Speed RatioRatio Geometry (mm/s) (%) (%) Flat 28 8.18 7.74 Flat 28 8.08 Flat 286.97 Flat 114 17.03 17.53 Flat 114 19.62 Flat 114 15.94 Flat 228 20.4421.70 Flat 228 21.24 Flat 228 23.41 6° Taper 28 7.87 8.32 6° Taper 288.77 6° Taper 114 19.84 18.48 6° Taper 114 16.55 6° Taper 114 19.04 7°Conical 28 8.37 10.56 7° Conical 28 12.05 7° Conical 28 11.25 7° Conical114 23.41 22.85 7° Conical 114 21.14 7° Conical 114 24.00 7° Conical 22821.71 21.37 7° Conical 228 19.50 7° Conical 228 22.91 Conical Flat 288.47 11.95 Conical Flat 28 13.25 Conical Flat 28 14.14 Conical Flat 11420.42 19.75 Conical Flat 114 19.22 Conical Flat 114 19.62 Conical Flat228 24.13 22.39 Conical Flat 228 23.31 Conical Flat 228 19.72

TABLE 41 Hole Expansion Ratio of Alloy 9 with Different Punch GeometriesMeasured Average Hole Hole Punch Expansion Expansion Punch Speed RatioRatio Geometry (mm/s) (%) (%) Flat 28 21.04 19.11 Flat 28 17.35 Flat 2818.94 Flat 114 24.80 24.29 Flat 114 19.74 Flat 114 28.34 Flat 228 26.0030.57 Flat 228 35.16 Flat 228 30.55 6° Taper 28 17.35 19.36 6° Taper 2819.06 6° Taper 28 21.66 6° Taper 114 29.64 31.14 6° Taper 114 32.14 6°Taper 114 31.64 7° Conical 28 22.63 24.05 7° Conical 28 23.61 7° Conical28 25.92 7° Conical 114 34.36 32.60 7° Conical 114 31.67 7° Conical 11431.77 7° Conical 228 36.28 36.44 7° Conical 228 38.87 7° Conical 22834.16 Conical Flat 28 27.72 25.59 Conical Flat 28 24.63 Conical Flat 2824.43 Conical Flat 114 30.28 32.64 Conical Flat 114 32.87 Conical Flat114 34.76 Conical Flat 228 32.90 35.45 Conical Flat 228 37.45 ConicalFlat 228 35.99

TABLE 42 Hole Expansion Ratio of Alloy 12 with Different PunchGeometries Measured Average Hole Hole Punch Expansion Expansion PunchSpeed Ratio Ratio Geometry (mm/s) (%) (%) Flat 28 29.38 31.59 Flat 2833.70 Flat 28 31.70 Flat 114 40.08 45.50 Flat 114 48.11 Flat 114 48.31Flat 228 50.00 49.36 Flat 228 40.56 Flat 228 57.51 6° Taper 28 29.9130.67 6° Taper 28 32.50 6° Taper 28 29.61 6° Taper 114 38.42 41.19 6°Taper 114 44.37 6° Taper 114 40.78 7° Conical 28 34.90 33.76 7° Conical28 33.00 7° Conical 28 33.37 7° Conical 114 45.72 49.10 7° Conical 11449.30 7° Conical 114 52.29 7° Conical 228 58.90 54.36 7° Conical 22853.43 7° Conical 228 50.75 Conical Flat 28 37.15 34.43 Conical Flat 2831.47 Conical Flat 28 34.66 Conical Flat 114 45.76 46.36 Conical Flat114 45.96 Conical Flat 114 47.36 Conical Flat 228 57.51 54.11 ConicalFlat 228 53.48 Conical Flat 228 51.34

This Case Example demonstrates that for all alloys tested, there is aneffect of punch geometry on edge formability. For all alloys tested, theconical punch shapes resulted in the largest hole expansion ratios,thereby demonstrating that modifying the punch geometry from a flatpunch to a conical punch shape reduces the damage within the materialdue to the punched edge and improves edge formability. The 7° conicalpunch geometry resulted in the greatest edge formability increaseoverall when compared to the flat punch geometry with the conical flatgeometry producing slightly lower hole expansion ratios across themajority of alloys tested. For Alloy 1 the effect of punch geometry isdiminished with increasing punching speed, with the three testedgeometries resulting in nearly equal edge formability as measured byhole expansion ratio (FIG. 79). Punch geometry, coupled with increasedpunch speeds have been demonstrated to greatly reduce residual damagefrom punching within the edge of the material, thereby improving edgeformability. With higher punch speed, the amount of heat generated atthe sheared edge is expected to increase and the localized temperaturespike may result in an annealing effect (i.e. in-situ annealing).Alternatively, with increasing punch speed, there may be a reducedamount of material transforming from the Recrystallized Modal Structure(i.e. Structure #4 in FIG. 1B) to the Refined High Strength NanomodalStructure (i.e. Structure #5 in FIG. 1B). Concurrently, the amount ofRefined High Strength Nanomodal Structure (i.e. Structure #5 in FIG. 1B)may be reduced due to the temperature spike enabling localizedrecrystallization (i.e. Mechanism #3 in FIG. 1B).

Case Example #17: HER in Commercial Steel Grades as a Function of HolePunching Speed

Hole expansion ratio testing was performed on commercial steel grades780, 980 and 1180. All specimens were tested in the as received(commercially processed) sheet condition.

Specimens for testing with a size of 89×89 mm were wire EDM cut from thesheet of each grade. The holes with 10 mm diameter were punched atdifferent speeds on two different machines with the same punch profilegeometry using a commercial punch press. The low speed punched holes(0.25 mm/s) were punched using an Instron 5985 Universal Testing Systemat 12% clearance and the high speed punched holes (28 mm/s, 114 mm/s,228 mm/s) were punched on a commercial punch press at 16% clearance. Allholes were punched using a flat punch geometry.

Hole expansion ratio (HER) testing was performed on the SP-225 hydraulicpress and consisted of slowly raising the conical punch that uniformlyexpanded the hole radially outward. A digital image camera system wasfocused on the conical punch and the edge of the hole was monitored forevidence of crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

Results from hole expansion tests are shown in Table 43 through Table 45and illustrated in FIG. 80. As it can be seen, the hole expansion ratiodoes not show improvement with increasing punching speed in all testedgrades.

TABLE 43 Hole Expansion Ratio of 780 Steel Grade at Different PunchSpeeds Punch Punch to die Sample Speed clearance Punch # (mm/s) (%)Geometry HER 1 5 mm/s 12% Flat 44.74 2 12% Flat 39.42 3 12% Flat 44.57 128 mm/s 16% Flat 35.22 2 16% Flat 28.4 3 16% Flat 36.38 1 114 mm/s 16%Flat 31.58 2 16% Flat 33.9 3 16% Flat 22.29 1 228 mm/s 16% Flat 31.08 216% Flat 31.85 3 16% Flat 31.31

TABLE 44 Hole Expansion Ratio of 980 Steel Grade at Different PunchSpeeds Punch Punch to die Sample Speed clearance Punch # (mm/s) (%)Geometry HER 1 5 mm/s 12% Flat 33.73 2 12% Flat 35.02 1 28 mm/s 16% Flat26.88 2 16% Flat 26.44 3 16% Flat 23.83 1 114 mm/s 16% Flat 26.81 2 16%Flat 30.56 3 16% Flat 29.24 1 228 mm/s 16% Flat 30.06 2 16% Flat 30.98 316% Flat 30.62

TABLE 45 Hole Expansion Ratio of 1180 Steel Grade at Different PunchSpeeds Punch Punch to die Sample Speed clearance Punch # (mm/s) (%)Geometry HER 1 5 mm/s 12% Flat 26.73 2 12% Flat 32.9 3 12% Flat 25.4 128 mm/s 16% Flat 35.32 2 16% Flat 32.11 3 16% Flat 36.37 1 114 mm/s 16%Flat 35.15 2 16% Flat 30.92 3 16% Flat 32.27 1 228 mm/s 16% Flat 27.25 216% Flat 23.98 3 16% Flat 31.18

This Case Example demonstrates that no edge performance effect based onhole punch speed is measurable in tested commercial steel gradesindicating that edge performance in conventional AHSS is not effected orimproved by punch speed as expected since the unique structures andmechanisms present in this application as for example in FIG. 1A andFIG. 1B are not present.

Case Example #18: Relationship of Post Uniform Elongation to HoleExpansion Ratio

Existing steel materials have been shown to exhibit a strong correlationof the measured hole expansion ratio and the material's post uniformelongation. The post uniform elongation of a material is defined as adifference between the total elongation of a sample during tensiletesting and the uniform elongation, typically at the ultimate tensilestrength during tensile testing. Uniaxial tensile testing and holeexpansion ratio testing were completed on Alloy 1 and Alloy 9 on thesheet material at approximately 1.2 mm thickness for comparison toexisting material correlations.

Slabs with thickness of 50 mm were laboratory cast of Alloy 1 and Alloy9 according to the atomic ratios provided in Table 2 and laboratoryprocessed by hot rolling, cold rolling annealing at 850° C. for 10 minas described in the Main Body section of this application.

Tensile specimens in the ASTM E8 geometry were prepared by wire EDM. Allsamples were tested in accordance with the standard testing proceduredescribed in the Main Body of this document. An average of the uniformelongation and total elongation for each alloy were used to calculatethe post uniform elongation. The average uniform elongation, averagetotal elongation, and calculated post uniform elongation for Alloy 1 andAlloy 9 are provided in Table 46.

Specimens for hole expansion ratio testing with a size of 89×89 mm werewire EDM cut from the sheet of Alloy 1 and Alloy 9. Holes of 10 mmdiameter were punched at 0.25 mm/s on an Instron 5985 Universal TestingSystem at 12% clearance. All holes were punched using a flat punchgeometry. These test parameters were selected as they are commonly usedby industry and academic professionals for hole expansion ratio testing.

Hole expansion ratio (HER) testing was performed on the SP-225 hydraulicpress and consisted of slowly raising the conical punch that uniformlyexpanded the hole radially outward. A digital image camera system wasfocused on the conical punch and the edge of the hole was monitored forevidence of crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

The measured hole expansion ratio values for Alloy 1 and Alloy 9 areprovided in Table 46.

TABLE 46 Uniaxial Tensile and Hole Expansion Data for Alloy 1 and Alloy9 Average Average Post Uniform Hole Uniform Total Elongation ExpansionElongation Elongation (ε_(pul)) Ratio Alloy (%) (%) (%) (%) Alloy 147.19 49.29 2.10 2.30 Alloy 9 50.83 56.99 6.16 2.83

Commercial reference data is shown for comparison in Table 47 from [PaulS. K., J Mater Eng Perform 2014; 23:3610.]. For commercial data, S. K.Paul's prediction states that the hole expansion ratio of a material isproportional to 7.5 times the post uniform elongation (See Equation 1).HER=7.5(ε_(pul))  Equation 1

TABLE 47 Reference Data from [Paul S.K., J Mater Eng Perform 2014;23:3610.] Post Uniform Hole Commercial Uniform Total ElongationExpansion Steel Elongation Elongation (ε_(pul)) Ratio Grade (%) (%) (%)(%) IF-Rephos 22 37.7 15.7 141.73 IF-Rephos 22.2 39.1 16.9 159.21 BH21019.3 37.8 18.5 151.96 BH300 16.5 29 12.5 66.63 DP 500 18.9 27.5 8.655.97 DP600 16.01 23.51 7.5 38.03 TRIP 590 22.933 31.533 8.6 68.4 TRIP600 19.3 27.3 8 39.98 TWIP940 64 66.4 2.4 39.1 HSLA 350 19.1 30 10.986.58 340 R 22.57 36.3 13.73 97.5

The Alloy 1 and Alloy 9 post uniform elongation and hole expansion ratioare plotted in FIG. 81 with the commercial alloy data and S. K. Paul'spredicted correlation. Note that the data for Alloy 1 and Alloy 9 do notfollow the predicted correlation line.

This Case Example demonstrates that for the steel alloys herein, thecorrelation between post uniform elongation and the hole expansion ratiodoes not follow that for commercial steel grades. The measured holeexpansion ratio for Alloy 1 and Alloy 9 is much smaller than thepredicted values based on correlation for existing commercial steelgrades indicating an effect of the unique structures and mechanisms arepresent in the steel alloys herein as for example shown in FIG. 1a andFIG. 1 b.

Case Example #19 HER Performance as a Function of Hole Expansion Speed

Slabs with thickness of 50 mm were laboratory cast from three selectedalloys according to the atomic ratios provided in Table 2 and laboratoryprocessed by hot rolling, cold rolling and annealing at 850° C. for 10min as described herein. Sheet from each alloy possessing theRecrystallized Modal Structure with final thickness of 1.2 mm were usedto demonstrate the effect of hole expansion speed on HER performance.

Specimens for testing with a size of 89×89 mm were cut via wire EDM fromthe sheet. Holes of 10 mm diameter were punched at a constant speed of228 mm/s on a commercial punch press. All holes were punched with a flatpunch geometry, and with approximately 16% punch to die clearance.

Hole expansion ratio (HER) testing was performed on an InterlakenTechnologies SP-225 hydraulic press and consisted of raising the conicalpunch that uniformly expanded the hole radially outward. Four holeexpansion speeds, synonymous with the conical ram travel speed, wereused; 5, 25, 50, and 100 mm/min. A digital image camera system wasfocused on the conical punch and the edge of the hole was monitored forevidence of crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

Hole expansion ratio values for the tests are shown in Table 48. Theaverage hole expansion ratio value is shown for each speed and alloytested showing an increase in HER values with increasing hole expansionspeed in all three alloys. The effect of hole expansion speed is alsodemonstrated in FIG. 82, FIG. 83, and FIG. 84 for Alloy 1, Alloy 9, andAlloy 12, respectively.

TABLE 48 Hole Expansion Ratio in Selected Alloys at Different ExpansionSpeeds Measured Average Hole Hole Hole Punch Expansion ExpansionExpansion Speed Speed Ratio Ratio Material (mm/s) (mm/min) (%) (%) Alloy1 228 5 19.09 20.55 22.54 20.02 25 30.70 28.58 29.14 25.91 50 34.0534.63 36.43 33.42 100 37.11 37.19 38.52 35.93 Alloy 9 228 5 34.06 34.1534.07 34.31 25 32.87 40.77 45.46 43.98 50 38.39 44.17 39.71 54.42 10048.01 49.50 55.27 45.23 Alloy 12 228 5 48.61 43.51 34.79 47.14 25 42.1350.64 57.82 51.96 50 63.77 62.97 68.46 56.68 100 57.79 56.73 49.28 63.11

This Case Example demonstrates that formability of the edge, i.e. itsability to be deformed with relatively reduced cracking, as measured byHER testing, can be affected by the speed of deformation of the holeedge (i.e. hole expansion speed). The alloys tested in this Case Exampledemonstrated a positive correlation between hole expansion ratio and thehole expansion speed, with increasing hole expansion speed resulting inrelatively higher measured hole expansion ratios.

Accordingly, in the broad context of the present disclosure, it has beenestablished that once an edge is formed, of any geometry by any edgeformation method which causes deformation of the metal alloy whenforming the edge (e.g. by punching, shearing, piercing, perforating,cutting, cropping, stamping), by increasing the speed at which that edgeonce formed is then expanded, one observes that the edge itself is thencapable of more expansion with a relatively reduced tendency to crack.The edge herein can therefore include an edge that defines an internalhole in a metal sheet of the alloys described herein, or an externaledge on such metal sheet. In addition, the edge herein may be formed ina progressive die stamping operation which is reference to metal workingoperation that typically includes punching, shearing, coining andbending. The edge herein may be present in a vehicle, or morespecifically, part of a vehicular frame, vehicular chassis, or vehiclepanel.

Reference to edge expansion herein is understood as increasing thelength of such edge with a corresponding change in the thickness of theedge. That is confirmed by the above data in Table 48, which shows thatwith respect to an edge that is present in a hole, when such edge in thehole is expanded at a speed of greater than or equal to 5 mm/min, oneobserves an increase in the hole expansion ratio (i.e. the edge in thehole is capable of expansion to higher percentages over the originaldiameter) and the edge getting thinner as shown for example in the crosssections of the expanded edges in FIG. 91.

Case Example 20 HER Performance as a Function of Punch Speed and HoleExpansion Speed

Sheet from Alloy 9 was produced according to the atomic ratios providedin Table 2. Slabs produced by continuous casting were hot rolled intohot band which was subsequently processed into sheet with thickness ofapproximately 1.4 mm by cold rolling and annealing cycles. Themicrostructure of the produced sheet using both SEM and etched opticalmicroscopy is demonstrated in FIG. 85 showing typical RecrystallizedModal Structure.

In FIG. 85A and FIG. 85B, the SEM micrographs shows the micron scalenature of the austenitic grains which contain some annealing twins andstacking faults. In FIG. 85C and FIG. 85D, etched samples were examinedusing optical microscopy. It can be seen that the grain boundaries arepreferentially etched and the microstructure showing the grainboundaries. The grain size was measured with the line intercept methodand is found to range from 6 μm to 22 μm with a mean value of 15 μm.

The sheet with Recrystallized Modal Structure was used for HER testing.Specimens for testing with a size of 89×89 mm were cut via wire EDM fromthe sheet. Holes of 10 mm diameter were punched at two different speedsof 5 mm/s using an Instron mechanical test frame and at 228 mm/s using acommercial punch press with a flat punch geometry and with punch to dieclearances of approximately 12.5% and 16%, respectively.

Hole expansion ratio (HER) testing was performed on an InterlakenTechnologies SP-225 hydraulic press and consisted of raising the conicalpunch that uniformly expanded the hole radially outward. Two holeexpansion speeds of 3 mm/min and 50 mm/min, synonymous with the conicalram travel speed, were used. A digital image camera system was focusedon the conical punch and the edge of the hole was monitored for evidenceof crack formation and propagation.

The initial diameter of the hole was measured twice with calipers,measurements were taken at 90° increments and averaged to get theinitial hole diameter. The conical punch was raised continuously until acrack was observed propagating through the specimen thickness. At thatpoint the test was stopped and the hole expansion ratio was calculatedas a percentage of the initial hole diameter measured before the startof the test. After expansion four diameter measurements were taken usingcalipers every 45° and averaged to account for any asymmetry of the holedue to cracking.

Hole expansion ratio values for tests are listed in Table 49. HER valuesvary from 2.4 to 18.5% in the samples with holes punched at 5 mm/s. Inthe case of 228 mm/s hole punching speed, HER values are significantlyhigher in a range from 33.8 to 75.0%. The effect of expansion speed isillustrated in FIG. 86. Increase in expansion speed results in higherHER values independent of utilized punching speeds (i.e. 5 mm/s and 228mm/s).

TABLE 49 Hole Expansion Ratio in Alloy 9 Sheet at Different Punching andExpansion Speeds Hole Hole Punch Punch Expansion Clearance Speed SpeedHER (%) (mm/s) (mm/min) (%) 16 228 3 33.8 16 228 3 41.3 16 228 50 63.116 228 50 75.0 12.5 5 3 2.4 12.5 5 3 7.9 12.5 5 50 12.7 12.5 5 50 18.5

The magnetic phases volume percent (Fe %) was measured in the HER testedsamples with different hole punching speed and hole expansion speedusing a Fischer Feritscope FMP30. The results are listed in Table 50.FIG. 87 illustrates the effect of on the magnetic phases volume percentin the tested samples as a function of the distance from the hole edge.As can be seen with higher punch speed and/or higher expansion speed,after testing is completed, the magnetic phase volume % increases nearthe hole edge and also away from the hole edge into the material. As theincrease in magnetic phase volume (Fe %) is consistent with increases inthe amount of Structure #5 in Table 1 which is formed duringdeformation, due to the formation of magnetic nanoscale alpha-iron fromthe starting non-magnetic austenite present in Structure #4.

TABLE 50 Magnetic Phases Volume (Fe %) in Alloy 9 at Different HolePunching Speeds and Hole Expansion Speeds as a Function of Distance fromHole Edge After Expansion Hole Creation and Expansion Parameters HolePunching Speed 228 228 228 228 5 5 5 5 (mm/s) Punch Clearance (%) 16 1616 16 12.5 12.5 12.5 12.5 Hole Expansion Speed 3 3 50 50 3 3 50 50(mm/min) HER (%) 33.8 41.3 63.1 75.0 2.4 7.9 12.7 18.5 Distance fromhole (mm) Magnetic Phases Volume % (Fe %) 1 27.3 31.6 37.9 39.1 7.1 9.513.5 20.1 2.5 17 21.1 29.6 36 2.4 2.7 6.5 6.2 4 6 7.5 17.4 24.6 0.94 1.12.4 2.4 5.5 2.2 2.8 6.3 11.3 0.47 0.45 0.96 0.75 7 0.82 0.89 2.8 4.40.21 0.29 0.38 0.28 8.5 0.33 0.35 1.3 1.9 0.23 0.22 0.24 0.16 10 0.210.21 0.66 1.1 0.21 0.2 0.2 0.13 11.5 0.15 0.16 0.42 0.67 0.2 0.18 0.210.12 13 0.13 0.14 0.26 0.37 0.18 0.18 0.22 0.11 14.5 0.12 0.13 0.25 0.310.19 0.18 0.23 0.11 16 0.13 0.14 0.31 0.38 0.19 0.19 0.22 0.13 17.5 0.20.22 0.53 0.84 0.19 0.2 0.24 0.14 19 0.16 0.25 0.37 0.61 0.2 0.22 0.220.12 22 0.11 0.13 0.21 0.24 0.19 0.21 0.22 0.1 25 0.12 0.12 0.19 0.230.19 0.2 0.2 0.11

This Case Example illustrates that the relative resistance to crackingof an edge as confirmed by HER testing can be increased by, in theexemplary case of forming an edge within a hole, by either increasinghole punching speeds, hole expansion speeds or both. The sheet fromAlloy 9, tested in this Case Example, demonstrated an increase in holeexpansion ratio with increasing hole punching speed (i.e. 5 to 228 mm/s)and/or the hole expansion speed (i.e. 3 to 50 mm/min). Accordingly,preferably herein for the subject alloys, one forms an edge in the alloyand expands the edge at a speed of greater than or equal to 5 mm/min,The magnetic phases volume percent (Fe %) in tested samples increaseswith increasing hole punching speed and/or the hole expansion speed overthe ranges studied. With this relatively greater amount of deformationavailable in and adjacent to the hole edge during the now disclosedincreased hole punching speed or hole expansion speed, the higher localformability and resistance to cracking of the edge is achieved in thematerial as measured by the HER.

Case Example #21 HER Performance as a Function of Hole PreparationMethod

Slabs with thickness of 50 mm were laboratory cast from three selectedalloys according to the atomic ratios provided in Table 2 and laboratoryprocessed by hot rolling, cold rolling and annealing at 850° C. for 10min as described herein. Sheet from each alloy possessing theRecrystallized Modal Structure with final thickness of 1.2 mm were usedto demonstrate an effect of hole expansion speed on HER performance.

Specimens for testing with a size of 89×89 mm were cut via wire EDM fromthe sheet. A 10 mm diameter hole was prepared by various methodsincluding punching, EDM cutting, milling, and laser cutting. Holepunching was done at a low quasistatic punching speed of 0.25 mm/s at16% punch to die clearance using a Komatsu OBS80-3 press. EDM cut holeswere first rough cut then the final cut was made at parameters to yielda visually smooth surface. During hole milling, holes were pilotdrilled, reamed to size, and then deburred. Laser cut samples were cuton a 4 kW fiber optic Mazak Optiplex 4020 Fiber II machine.

Hole expansion ratio (HER) testing was performed on an InterlakenTechnologies SP-225 hydraulic press and consisted of raising the conicalpunch that uniformly expanded the hole radially outward. In FIG. 88, theresults of HER testing is provided for each alloy as a function of thehole preparation method. As shown, in the case of punched holes, HERvalues are the lowest for all three alloys in a range from 6 to 12%.Samples with EDM cut, milled and laser cut holes exhibit high HER valuesfrom 65 to 140%+. Note that the ˜140% expansion represented the maximumextension limit of the press crosshead during testing so, in the sampleswith EDM cut holes from Alloy 12, and with milled holes from Alloy 9 andAlloy 12, the expansion limit was not reached during HER testing (i.e.actual value>140%).

In FIG. 89, SEM images of the sample cross section near the hole edgeprior to expansion are provided at low magnification for samples fromAlloy 1 with holes prepared by different methods. In the punched sample(FIG. 89A), the typical rollover zone at the top and burr zone at thebottom can be seen. Additionally, a hemispherical shear affected zone isvisible at the edge of the hole with penetration of ˜0.5 mm at thedeepest point. A similarly shear affected zone was observed in punchedsamples from the other two alloys as well but not in any of the sampleswith holes produced by the non-punching methods. Note that every methodutilized for hole preparation introduced some kind of defects at thehole edge. In the EDM cut hole (FIG. 89B), the edge is perpendicular ona cross section image but small micron scaled cutting defects can beseen at the surface; in the milled samples (FIG. 89C) the edge of theholes is trapezoidal in shape; and in the laser cut holes (FIG. 89D),the edge wandered in a sideways fashion as the laser penetrated thesample. In FIG. 90, SEM images of the cross sections near the hole edge(i.e. at the edge and up to 0.7 mm away from the edge) prior toexpansion are provided at higher magnification for samples from Alloy 1with holes prepared by different methods including punching at a holepunching speed of 0.25 mm/s, EDM cut hole, milled hole, and laser cuthole. The microstructure near the hole edges are illustrated in FIGS.90A, 90B, 90C and 90D respectively. As can be seen, the edge of the holepunched at 0.25 mm/s (FIG. 90a ) is relatively highly deformed therebyleading to the observed low HER values. This structure near the edge ofthe punched sample is representative of Structure #5 Refined HighStrength Nanomodal Structure in Table 1 whereby the structures near thehole edge of the EDM cut, milled, and laser cut holes, is representativeof Structure #4 Recrystallized Modal Structure in Table 1. However, inexamples where holes were prepared by non-punching methods (FIGS. 90B,90C, 90D), the resulting alloys experienced excellent local formabilitywith high HER values from 65 to 140%+ consistent with the ductile natureof Structure #4 near the hole edges. In FIG. 91A (punched hole), 91B(EDM cut hole), 92C (milled hole), and 91D (laser cut hole), SEM imagesof the cross section near the hole edge after HER testing are providedat low magnification for samples from Alloy 1. Note that the thicknessof samples near the hole is smaller in the expanded samples with higherHER values since the expansion of the holes results in sample thinningnear the hole edge.

In FIG. 92, images of sample cross sections near the hole edge after HERtesting (i.e. after expansion until failure by cracking) are provided athigher magnification for samples from Alloy 1 with holes prepared bydifferent methods showing similar deformed structure in all cases. Sincehole expansion and deformation of the edge is complete, themicrostructure near all of the hole edges are similar and representativeof Structure #5 Refined High Strength Nanomodal Structure in Table 1.This Case Example demonstrates the effect of edge preparation on theresulting local formability in alloys herein. Punching at a low speedsof 0.25 mm/s is causing structural changes near the hole edge consistentwith previous case examples resulting in limited local formability ofthe edges and low HER values. However, in examples where holes wereprepared by non-punching methods, the resulting alloys experiencedexcellent local formability with high HER values from 65 to 140%+consistent with the ductile nature of the microstructure in the samplesand at the hole edges.

What is claimed is:
 1. A method for expanding the edge of an alloycomprising: a. supplying a metal alloy comprising at least 50 atomic %iron and at least four elements selected from Si, Mn, B, Cr, Ni, Cu orand melting said alloy and cooling at a rate of <250 K/s or solidifyingto a thickness of >2.0 mm up to 500 mm and forming an alloy having a Tm;b. heating said alloy to a temperature in a range of 700° C. to belowsaid Tm and reducing said thickness of said alloy at a strain rate of10⁻⁶ to 10⁴ to provide a first resulting alloy having an ultimatetensile strength of 921 MPa to 1413 MPa and an elongation of 12.0% to77.7%; c. stressing said first resulting alloy and providing a secondresulting alloy having an ultimate tensile strength of 1356 MPa to 1831MPa and an elongation of 1.6% to 32.8%; d. heating said second resultingalloy to a temperature below said Tm and forming a third resulting alloyhaving an elongation of 6.6% to 86.7%; e. forming an edge in said thirdresulting alloy and expanding said edge at a speed of greater than orequal to 5 mm/min to form an expanded edge.
 2. The method of claim 1wherein said edge is expanded at a rate in the range of greater than orequal to 5 mm/min to 100 mm/min.
 3. The method of claim 1 wherein saidalloy comprises Fe and at least five elements selected from Si, Mn, B,Cr, Bi, Cu or C.
 4. The method of claim 1 wherein said alloy comprisesFe and at least six elements selected from Si, Mn, B, Cr, Ni, Cu or C.5. The method of claim 1 wherein said alloy comprises Fe, Si, Mn, B, Cr,Ni, Cu and C.
 6. The method of claim 1 wherein said third resultingalloy has a yield strength ranging from 197 to 1372 MPa.
 7. The methodof claim 1 wherein said third resulting alloy has an ultimate tensilestrength ranging from 799 to 1683 MPa.
 8. The method of claim 1 whereinbefore step (e) the edge in said alloy is exposed to a temperature in arange of of 400° C. to below said Tm.
 9. The method of claim 1 whereinin step (e), the edge defines either an internal hole and/or an externaledge.
 10. The method of claim 1 wherein in step (e), the edge is formedby punching, piercing, perforating, cutting, cropping, EDM cutting,waterjet cutting, laser cutting, or milling.
 11. The method of claim 1,wherein said edge is formed in a progressive die stamping operation. 12.The method of claim 1, wherein said expanded edge is positioned in avehicle.
 13. The method of claim 1, wherein said expanded edge is partof a vehicular frame, vehicular chassis, or vehicular panel.